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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 13
45/2 March/April 2009
Editor's Note PM Industry News in Review Company Profile Magnesium Elektron Powders PMT Spotlight On …Stan Cuthbert Consultants’ Corner Joseph Tunick Strauss
FOCUS: PM Metallography 17 Cutting-Edge PM Metallography T.F. Murphy, FAPMI
19 Three-Dimensional Characterization and Modeling of Porosity in PM Steels N. Chawla, J.J. Williams, X. Deng, C. McClimon, L. Hunter and S.H. Lau
29 Characterization of Powders and PM Components Utilizing Transmission Electron Microscopy C. Blais, G. L’Espérance, FAPMI, and P. Plamondon
39 Porosity Statistics and Fatigue Strength of Sintered Iron P. Beiss, FAPMI, and S. Lindlohr
49 Evaluation of PM Fracture Surfaces Using Quantitative Fractography T.F. Murphy, FAPMI
DEPARTMENTS 62 Meetings and Conferences 63 PM Bookshelf 64 Advertisers’ Index Cover: Epoxy skeleton showing morphology of pore network inside a compact. Photo courtesy Thomas Murphy, Hoeganaes Corporation.
The International Journal of Powder Metallurgy (ISSN No. 0888-7462) is a professional publication serving the scientific and technological needs and interests of the powder metallurgist and the metal powder producing and consuming industries. Advertising carried in the Journal is selected so as to meet these needs and interests. Unrelated advertising cannot be accepted. Published bimonthly by APMI International, 105 College Road East, Princeton, N.J. 08540-6692 USA. Telephone (609) 4527700. Periodical postage paid at Princeton, New Jersey, and at additional mailing offices. Copyright © 2009 by APMI International. Subscription rates to non-members; USA, Canada and Mexico: $100.00 individuals, $230.00 institutions; overseas: additional $40.00 postage; single issues $55.00. Printed in USA 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:
[email protected]
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EDITOR’S NOTE
T
he ability to study surfaces and to delve into the interior of solids at a resolution approaching the atomic level serves as the basis for our understanding of material properties, in particular the behavior of solids under stress. To this end, metallography is the predominant and most powerful technique available to the materials scientist. This Focus Issue, coordinated by Tom Murphy, surveys cutting-edge metallographic techniques that characterize the complexities of the internal and external microstructures of powder metallurgy (PM) materials in which porosity is an intrinsic component. In addition to its analytical capability, metallography frequently exposes the aesthetic qualities of a microstructure. This duality is clearly illustrated by the PM micrograph on the front cover—an epoxy skeleton showing the morphology of the pore network that remains after extraction of the iron from a sintered compact using dilute acids. Joe Strauss again brings his unique flair and expertise to the “Consultants’ Corner.” The issues he addresses are quality assurance in PM part fabrication, the role of PM in rapid prototyping and the choice of powder grades, and a comparison of cold-spray and thermal-spray processing. Amid the current and widespread economic doom and gloom, Peter Johnson’s “Company Profile” on Magnesium Elektron Powders should evoke a welcome sense of confidence for the future. The company has carved out an important niche in the specialty metal powders market by focusing on advances in technology, quality, product customization, and new product development, as well as a rapid response to customer needs.
Alan Lawley Editor-in-Chief
Now in its second year, the e-version of the Journal, a facsimile of the familiar hard copy, offers APMI members user-friendly navigational features and powerful search capabilities. Of particular convenience is the ability of the software to locate technical information (articles, references, etc.) from previous issues of the Journal. To log in, all that is needed is a personal user ID and password. For no particulate reason, I recently found myself cleaning our old files. On reaching the letter “E”, I rediscovered a dog-eared file titled “English Food and Recipes.” The contents of this large file brought back instant memories of a misspent youth in post–World War II England. Taking license from the fact the PM plays an important role in the processing of food and beverages, I thought it might be of interest to share with you the essence of the file in the form of a quiz! Can you identify the primary constituent(s) of the following delicacies from the United Kingdom? Better still, have you ever had the privilege of following traditional recipes to prepare these dishes? Bangers and Mash; Black Pudding; Bubble & Squeak; Cornish Pasties; Haggis; Pikelets; Pork Pie; Spotted Dick; Toad-in-theHole; Yorkshire Pudding. Bon appetit!
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Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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PM INDUSTRY NEWS IN REVIEW
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.”
New ISO Certification The production and distribution operations of Advanced Metalworking Practices, LLC (AMP), Carmel, Ind., have been certified to ISO 9001:2000 (Without Design) International Quality System standard. AMP makes feedstocks for the metal injection molding (MIM) industry. New Tungsten Plan The Defense National Stockpile Center, Fort Belvoir, Va., has changed its awards announcement policy. It will report tungsten awards twice in a fiscal year, in April to announce first half sales and in October to announce second half sales. GKN Updates Performance GKN plc expects group revenues for 2008 will rise by about 12 percent and profit before taxes will approach £150–£170 million (about $215–$243 million). However, the company reports that conditions in global automotive markets have continued to decline since its last update. PowderMet2009 HOTEL OFFERS UNPRECEDENTED RATES In light of current economic and market conditions, MPIF has finalized a new agreement with The Mirage Hotel in Las Vegas, headquarters for PowderMet2009 this June, that will help maximize attendance. Reduced MPIF confer-
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ence rates are now only $95.00 per night, including weekends. Visit www.mpif.org for details. Industry Loses Noted Powder Metallurgist Howard I. Sanderow, 64, died on February 2. He was president of Management & Engineering Technologies (MET Group Inc.), a PM consulting company he founded with his wife Barbara in 1988. Powder Maker Reports Weak Fourth Quarter Höganäs AB, Sweden, production declined 28 percent in the fourth quarter of 2008 because of sharply reduced demand in all market regions, including Asia and South America. Net sales for the full year increased 4.5 percent to MSEK 6,103 (about $740 million), mainly attributed to price increases and higher scrap metal surcharges. President Obama Learns About PM President Barack Obama showed a serious interest in powder metallurgy during his visit to CMW Inc., Indianapolis, Ind., in 2008, reports Mark B. Gramelspacher, president & CEO. Obama spent several hours at CMW, a manufacturer of PM electrical contacts, PM tungsten alloys, and resistance welding consumables, in April 2008 while on a campaign swing through the Midwest.
PRIMA Collaborates in Roller-Contact Fatigue Testing Prima Business Specialists LLC, State College, Pa., and Product & Assurance & Services, Ridgway, Pa., will collaborate to provide Hoffman Roller Contact Fatigue (RCF) tests and analysis to prospective users. The RCF test provides early detection and remedial action for potentially damaging cracks during the manufacturing process. PowderMet2009 Las Vegas Program Now Online The complete program as well as registration and hotel information for PowderMet2009 in Las Vegas are now online at www.mpif.org/ meetings/2009/09_gateway.htm. Sponsored by the Metal Powder Industries Federation and APMI International, the PM industry’s largest annual conference will take place in Las Vegas, Nevada, June 28–July 1. Former Metaldyne Employees Sentenced Three former employees of Metaldyne Corp. were sentenced to prison terms in federal court in Detroit on February 13. They pleaded guilty to charges in connection with a conspiracy to steal confidential information about the company’s powderforged connecting rod manufacturing costs and processes. ijpm
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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PM INDUSTRY NEWS IN REVIEW
PM Parts Maker Adjusts Capacity Miba AG, Laakirchen, Austria, is addressing the crisis of declining demand from the international automotive market by adjusting capacity. The company’s measures include reducing employee overtime and excess hours, using up residual holidays, and cutting temporary workers and educational leave. Automotive Suppliers Face Challenging Business Conditions International automotive suppliers Tomkins plc and GKN plc report declining profits in 2008. Business conditions for both deteriorated especially during the final quarter.
Tungsten Sales Rise North American Tungsten Corp. Ltd. (NTC), Vancouver, B.C., Canada, reports a 49.7 percent sales increase to Can$17.6 million for the first fiscal quarter of 2009, ending December 31, 2008. NTC’s Cantung mine produced 79,978 metric ton units of tungsten concentrate, a 22 percent increase over the comparable quarter the previous fiscal year. Tungsten Powder Expansion Shelved H.C. Starck Canada Inc. has suspended further investment in expanding capacity at its Sarnia, Ontario, tungsten powder plant. The company attributed the decision to the cooling economic
climate and the resulting decline in demand in the North American tungsten market. New Developments to Be Aired at Refractory Metals and Hard Materials Meeting The Plansee Group announced the program of the 17th Plansee Seminar at its headquarters in Reutte, Austria, May 25–29. More than 200 presentations will focus on PM processing of refractory metals, hard materials and composite materials, including applications in energy, information technology, communications, and lighting. 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 2, 2009 International Journal of Powder Metallurgy
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2009 International Conference on Powder Metallurgy & Particulate Materials June 28–July 1, The Mirage Hotel, Las Vegas
• International Technical Program • Worldwide Trade Exhibition • Special Events
For complete program and registration information contact: INTERNATIONAL
METAL POWDER INDUSTRIES FEDERATION APMI 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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COMPANY PROFILE
Magnesium Elektron Powders By Peter K. Johnson* Magnesium Elektron Powders (MEP) services customers all over the world from plants in New Jersey, Pennsylvania, and Ontario, Canada. “We make powders, chips and granules for many markets, as well as market wrought and cast alloys,” says James G. Gardella, president. “Our Reade Manufacturing Co. in Manchester, New Jersey. has been a magnesium powder supplier since 1941.” The company is part of Magnesium Elektron, a division of the Luxfer Group. Luxfer employs 2,000 workers in 21 plants worldwide. In addition to magnesium powders and granules, Luxfer supplies high-pressure specialty gas cylinders, zirconium chemicals, aluminum and composite parts, wrought and cast magnesium alloys, and magnesium sheets & plates. Broadly speaking, magnesium’s most important uses are as an alloying agent in aluminum production, as die casting alloys, and as magnesium powders. MEP consists of three companies with a total of 100 employees: Reade Manufacturing Co., Hart Metals Inc., and Niagara Metallurgical Products, Ltd. Reade was acquired by Magnesium Elektron in 1990. Located on a 250,000 sq. m (62-acre) site (Figure 1), Reade makes magnesium and specialty alloy powders by mechanical comminution, better known as grinding. Magnesium ingot is converted into chips, granules, and coarse or fine powder using the following steps: chipping, grinding (coarse, fine, and ultrafine), screening, and blending. Specialty alloys are processed by crushing, grinding, and screening. A new automated production line for specialty alloys was recently added at the facility, Figure 2. Acquired in 1998, Niagara Metallurgical Products, Ltd. operates from a 1,340 sq. m (14,400 sq. ft.) plant in Stoney Creek, Ontario, Canada, Figure 3. In addition to producing magnesium *Contributing editor and consultant
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
Figure 1. Reade Manufacturing Co., Manchester, New Jersey
Figure 2. New specialty alloy production line at Reade
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COMPANY PROFILE: MAGNESIUM ELEKTRON POWDERS
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chips and powders using mechanical grinding, the plant also has a blending capability. Hart Metals Inc. atomizes pure magnesium spherical powder via a proprietary technology. The company was acquired by MEP in 1998. Located on a 28,200 sq. m (7-acre) site in Tamaqua, Pennsylvania, the plant is undergoing a major expansion that will more than double its magnesium atomization capacity, Figure 4. “All three of our plants use chipping systems and mechanical grinding for magnesium chips, granules, and powders as well as screening and blending,” reports John E. McConaghie, vice president–operations. MEP atomized magnesium powders are available in a particle-size range from 5 µm to 5 mm for various applications and blending with aluminum PM grade powders. Specialty metal and alloy powders (Fe-V, Fe-Al, Fe-B, Fe-Nb, and Fe-Si) are made for welding applications like stick electrodes and cored-wire. The company expects to add additional specialty alloy powders to their product portfolio. Quality and safety are critical in MEP’s production processes, notes Deepak Madan, vice president, technology & new product development. “We
adhere to stringent military specifications as well as ISO 9001:2000 and ISO 14001:2004 registrations,” he says. “And all of our materials undergo 100 percent testing and sampling.” The in-house laboratory at the Manchester facility is typical of MEP’s other plants, Figure 5. Powder characterization includes Rotap sieve analysis, Microtrac laser diffraction, and a sedigraph analyzer. Careful documentation of customer orders provides details on materials purchased decades ago. MEP follows strict safety precautions in the handling of its materials and adheres to National Fire Protection Association guidelines. Production units are separated and powder handling is minimized. Magnesium chips, granules and powders run a wide gamut of applications, from defense and aerospace to pharmaceutical, industrial, and consumer products. The defense market accounts for the most interesting usage of magnesium powder. Typical applications are aircraft infrared counter -measure flares, illuminating and marker flares, and tracer bullets, Figure 6. Improvised explosive device (IED) simulators used in military training are a new
Figure 3. Niagara Metallurgical Products, Ltd.
Figure 4. Hart Metals Inc. atomizing plant
Figure 5. In-house laboratory monitors quality control
Figure 6. Military aircraft counter-measure flares use magnesium
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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COMPANY PROFILE: MAGNESIUM ELEKTRON POWDERS
application using magnesium. MEP’s high-purity magnesium chips and granules are used in the chemical and pharmaceutical industries. Magnesium granules and powders are also used for the manufacture of pesticides, herbicides, fungicides, chemical synthesis, and for the desulphurization of steel. Special-effect pyrotechnics use magnesium metal and alloy powders. Magnesium powders are mixed with aluminum powder to create PM pre-blends. Specialty alloy powders are used as sintering aids and alloying additives in PM. There is also a new interest in developing lightweight parts for military and automotive applications. Flameless ration heater (FRH) pads represent a growing market for Magnesium Elektron Powders. FRH pads consist of a proprietary blend of a magnesium-alloy with other ingredients. When acti-
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
vated with water, these pads generate heat without the need for any exter nal power source. Applications include self-heating ready-to-eat meals (MRE) for the military and consumers. Reusable packs that provide hot water on demand are a new application that will be available to consumers in the near future. For example, 1.5 L of water can be heated in 20 to 30 min for coffee, tea, or hot chocolate. Though relatively unknown in the conventional PM parts business, Magnesium Elektron Powders has carved out an important niche in the specialty metal powders market. “Our strengths are keeping technology current, high-quality materials, product customization, new product development, and serving customer needs quickly,” stresses James Gardella. “We are confident about our future growth.” ijpm
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SPOTLIGHT ON ...
STAN CUTHBERT, PMT Education: BS, Metallurgical Engineering, Michigan Technological University, 1969 Why did you study powder metallurgy/particulate materials? My interest in powder metallurgy (PM) goes back a generation to my father (also Stan Cuthbert). He worked at RB&W, a PM facility in Coldwater, Michigan. Occasionally he would take me to work with him on Saturdays. He must have been ahead of his time because it was long before “take a child to work” was ever conceived. I was fascinated by the pounding of the presses and especially the sintering furnaces. I knew then I wanted to understand how a part that can crumble in your hand could pass through a sintering furnace and come out strong. When did your interest in engineering/ science begin? It is amazing how certain individuals can shape your life. I can trace my interest in science back to my sixth grade teacher, Mr. Sistanich. It was the year of the first Sputnik launch and he was most encouraging and supportive of my interest, helping me obtain information on the U.S. and Russian space programs. What was your first job in PM? What did you do? My first job happened by accident. I was interviewed by Glidden Metals (SCM Metals). Glidden was expecting a minerals beneficiation (extractive) metallurgist to work on alternative methods of producing copper powder, but my degree was in physical metallurgy. As an afterthought I was introduced to Arthur Backensto (APMI International Fellow, 2003) and was hired for his research group. There I began my career developing brazing and infiltrating materials. At that time interest was developing in a brazing material for the PM rotor in the Wankel engine. At Glidden, as sole inventor, I was awarded patents for an infiltrating paste and for brazing and soldering compositions. Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
Describe your career path and companies worked for, and responsibilities. I moved within the company to the SCM Metals group in Johnstown, Pennsylvania. There I was responsible for bringing their dispersion-strengthened copper materials from the research group into production. Today the “Glidcop” materials are probably best known for their use in the welding electrode industry. However, in one of the early commercial uses, this high-strength, high-conductivity material was extruded into bars and used as heat sinks for the welding of the Alaskan pipeline. My next job was at Vickers Hydraulics where I thought I may be moving away from PM but, instead, I began working with PM parts to supply hardware for the hydraulics industry. This included hardware from large Ferro-Tic rotors to small copper-infiltrated piston shoes and low-carbon iron solenoids. The current stop in my career is at TRW, an independent supplier to the global automotive industry, and a leader in active and passive safety technology. The applications that we find for PM continue to grow. The variety of materials used keeps the work interesting, embracing everything from standard iron–carbon to MIM stainless steel. Our suppliers employ conventional sintering, sinter hardening, and vacuum sintering. What gives you the most satisfaction in your career? Without a doubt it is the variety of challenges and the fact that I have never strayed far from basic metallurgical laboratory work. To me, there is nothing more exciting than working on a failure analysis problem involving a combination of scientific knowledge, detective work, Metallurgist TRW Inc. 4505 West 26 Mile Road Washington, Michigan 48094 Phone: 586-786-7640 E-mail:
[email protected]
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SPOTLIGHT ON ...STAN CUTHBERT, PMT
common sense—and almost always under pressure to complete the project immediately! List your MPIF/APMI activities. I am a Past Chairman of the Michigan chapter of APMI International. I was also a speaker at APMI meetings, giving a presentation on “Powder Metal’s Role as a Substitute for Strategic Materials.” What major changes/trend(s) in the PM industry have you seen? The PM industry has previously described itself, in its own literature, as a “black art.” This concept has given way to automated handling systems, dedicated cells, and fully computer-monitored sintering and heat-treating furnaces. Now and in the future the emphasis will remain on global sourcing from a parts-purchasing point of view. I now have to be familiar with MPIF Standard 35, and also global material specification systems. In the automotive world the demand for quality and lower pricing will continue to increase. That is where PM will continue to shine with improvements in techniques for producing near -net shapes at higher output rates and with enhanced tool life. Why did you choose to pursue PMT certification? I had been out of the mainstream PM industry for awhile so it was natural for me to want to know if I had kept up with the technologies. Pursuit of certification was a natural way to find out. In addition, I was continuing to be increasingly involved in PM part selection within TRW.
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How have you benefited from PMT certification in your career? I was surprised at how well APMI distributes its technical information, and by how many people use it. There have been instances where our hardware engineers have been in casual discussions with PM parts suppliers and the suppliers have asked to have me join the discussion because they knew I was here. Maybe they did some homework by checking the Who’s Who in PM and observed that I had PMT Level I certification! This has happened more than once. What are your current interests, hobbies, and activities outside of work? Again, let me go back to my father, who is an avid golfer. He told me if I wanted to live a long, stressfree life, “do not take up golf.” So unlike many of your readers, I do not play golf. Playing golf in college presented a challenge in the snowbound upper peninsula of Michigan—but at least I met my wife Cookie there. Several years ago I realized that in my entire career I have only written technical reports, so I started taking classes in creative writing. In my spare time I now write poetry. I am currently compiling my first book, tentatively titled, “A Stain on Stainless Steel.” Look for it in the future! ijpm
Would you like to be featured here? Have you been PMT Certified for more than 2 years? Contact Dora Schember (
[email protected]) for more information.
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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CONSULTANTS’ CORNER
JOSEPH TUNICK STRAUSS* Q
Our PM process periodically produces parts out of specification (low density or out of geometric tolerance). We have tightened up on our process and our Cpk is capable of 1.33, which we have promised to our customer or we lose the order. We have calibrated all pertinent transducers (temperature, pressure, flow rates) and checked the validity of the incoming material certifications. Regardless, we still have the same periodic problem. What should we do? Let me preface a response by recounting a story that Matt Bulger (NetShape Technologies—MIM) passed on to me. A Fortune 500 company rolls into a small PM house and demands of the QA manager, "I need all critical dimensions on the parts you make for us to have a 2.00 Cpk or better! Can you make this happen?" The QA manager says, "Sure, if you'll grant me one small request." Fortune 500 company: "I guess so, what is it?" QA manager: "Change all nominal dimensions and tolerances to those that the tooling creates and which my process produces." If it were only so easy! But do not dismiss this because it is fantasy. In many instances this discussion needs to take place, if for no other reason than to point out that there are two sides to any specification issue—the producer (striving to meet the specification), and the consumer (setting what the specification should be). In this case you state that the system is “capable” of producing a Cpk of 1.33 and it is this level of control that is required by the customer. Are you implying that you do not achieve this level of control all of the time, but you bid on the job as if you could? There are two issues: 1. What can be done to improve the process control of your manufacturing process? 2. What can be done to achieve a specification that is compatible to you and the customer?
A
Issue 1: It is not unusual to see companies that either inherited a process or developed a process and then fired all the cognizant engineers. The facility is now run by equipment “operators," expediters on the floor, and some managers up front. There is not an engineer in sight. Or, if there is an engineer, he/she is too busy adhering to ISO, putting out fires, or solving banal scheduling problems to examine the problem. This is what keeps me employed as a consultant! To companies that fire their engineers, or do not allow their engineers to do engineering, thank you. It is encouraging to note that your process is within specification, your equipment works, and your incoming material meets specification. This tells me that the entire process was set up on the edge of workability. Has there been any sensitivity analysis done on any of the inputs? A process should be designed with as high a margin of safety around each input as possible; this is what makes a robust process. It is easy to design a process that falls off precipitously under the right (or wrong) combination of conditions. This is one of the dangers of blindly following the statistics as an “end-all” for characterizing a process, especially when the distribution of the process outputs may not be centered on the mean. Returning to your specific problem, notwithstanding all the statistical and quality controls you have put in place, these are just measurements and these tools cannot replace a fundamental understanding of the process. There is a metallurgical and/or mechanical reason that your parts are out of specification, regardless of what statistical tests of the process tell you. And, a Cpk of 1.33
*Engineer & President, HJE Company, Inc., 820 Quaker Road, Queensbury, New York 12804; Phone: 518-792-8733, Fax: 518792-8735; E-mail:
[email protected]
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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CONSULTANTS’ CORNER
today will be a Cpk of 1.42 tomorrow, 1.21 the next day, and 1.19 the day after that, etc. If you are close to a cliff, you will fall off eventually, and the cause may be a small number of trivial changes that combine to make the “perfect storm.” Get the help you can from suppliers, consultants, or from your workforce. Work the problem to the best of your ability until you get that margin of error. Issue 2: Here is a paraphrased quote from Arlan Clayton, the first recipient of the Kempton H. Roll Powder Metallurgy (PM) Lifetime Achievement Award and an individual who has been around the PM block a time or two: "90% of all quality issues are really business issues." So, with this in mind ask yourself: “Why did I win this contract? Did I win it on price? Why was my price lower than the competition? Is it because I had a fundamental economic advantage or because everyone else quoted the part with a secondary operation that they estimated was needed to bring the part into specification, and I missed it?” Many quality issues would pose no problem at all if you can get the price you need; the price that makes you profitable. Then, when there are rejectable parts, it does not throw the project immediately into the red. You should be making a profit even with some finite reject rate. If a secondary operation is needed to achieve the quality level, this is satisfactory as long as it is accounted for in the quote, and not added on after the customer has accepted your bid. When the basic process capability intersects with a low-margin part, the trouble begins. No manufacturing process is perfect, so do not bid on this assumption, especially when you do not have one. If you have a good process, then bid with your good process. In any case, bid with the process you have, not the one you want, or the customer demands. You may have a perfectly good process at a Cpk of 1.2 but it becomes a failure when filling an order that requires a higher Cpk. Improving your process under the gun is always more expensive than normal evolutionary improvement. What to do? The first step is to be honest with the customer. The cold hard fact is that if you cannot solve this issue, sooner or later you will do one of three things: (i) drop the part, leaving the customer in the lurch with no supply on a custom PM part; (ii) raise the price, but that may make the customer's end product uncompetitive and kill the
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program; or (iii) start cutting inspection corners, leading to line problems with the customer, retur ned shipments, and messy arguments. Sooner or later you will do either (i) or (ii) anyway. Customers do not like to hear about problems from suppliers, but they know problems develop. The issue is how to deal with them. If a supply chain is threatened, or a dramatic price increase looms, you will be amazed at how quickly a tolerance may be loosened, or how a measurement technique may be changed to bring a part into specification. But this should be done up front, not after the fact. The discussion has to be initiated by the supplier. Unless you tell the customer there is a problem, how will he/she know? There are always people in your company unwilling to admit a mistake, or unwilling to take any "blame" in a project. In the end, the facts will be the facts, and oftentimes customers will prefer an honest supplier to one that never comes clean until the dam breaks. Give yourself and your customer a chance to find a solution. Getting into this kind of predicament with a PM parts customer is bad for your company. However, you must also realize that getting into a predicament with a first-time application of PM is bad for the entire PM industry. What is the role of metal powder in rapid prototyping (RP)? Is there a grade of powder that RP requires? RP and rapid manufacturing (RM) continue to undergo considerable development and their applications and markets are expanding rapidly. Many of the RP and RM technologies involve the use of metal powder, and although the volume used may never compete with those for press-andsinter or metal injection molding (MIM), RP and RM should be embraced by the powder metallurgy (PM) industry. RP and RM are manufacturing technologies that build the part by additive methods and also without specific or dedicated tooling. That is, rather than starting with a bulk material and cutting away excess material until you have a part, RP and RM build the part using essentially only the material used in the final part. In addition, RP is a totally flexible process in which all parts can be made using the same basic tools or method; no fixed tooling is necessary. The manufacturing strategy in RP is usually to divide the part into
Q A
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planes of finite thickness and to build up the part layer by layer. The use of CAD/CAM is an integral and essential part of rapid prototyping. It should be noted that there is a difference in RP and RM. RP is the manufacture of a prototype part, which implies that the part is a compromise in the actual part desired. For instance, a metal part may be prototyped in a plastic but obviously the part, e.g., a turbocharged rotor, cannot be tested in service. The RP’d part provides the shape but not necessarily all of the attributes of the desired part. RM is the production of a part to be used in service. Thus, the material used in RM must be equivalent to the material used in other manufacturing methods. In many cases plastics are adequate. However, the holy grail of RP and RM is to be able to directly make metal parts. In order to build a part layer by layer, the material must be applied layer by layer. For RP/RM methods that use powder as the material it is necessary to introduce a layer of powder of a precise thickness. Once the powder is distributed in a uniform layer the portions of the layer that are to be incorporated into the part are “fixed” in place. In general, two methods are used: (1) application of a binder, and (2) laser “sintering.” In the first case a binder (generally organic) is applied using technology similar to the printhead in your inkjet printer, which allows precise placement of the binder. The final three-dimensional (3D) part is similar to a MIM part in that it must be debound and sintered to obtain the attributes of a fully metallic part. In the case of laser sintering, the laser bonds the particles by melting the powder. This can entail full or partial melting or liquid-phase sintering and high densities are achievable. Machines using an electron beam for melting/sintering are also available. Regardless of the RP/RM method used, one of the most important properties of the powder is its ability to be able to be doctored into a uniformly dense layer of a precise thickness. Certainly the flow properties of the powder are important, which implies that coarse spherical powder is desired. However, coarse material will be difficult to sinter to high densities, especially for processes using the binder print method. Powder with a maximum packing factor is also desired to minimize shrinkage, which allows for tighter control of dimensional tolerance. For laser or electron-beam sintering, a high packing factor is also desirable to provide adequate ther mal mass to prevent “punch through” or excessive melting. However, high packVolume 45, Issue 2, 2009 International Journal of Powder Metallurgy
ing factors and flowability do not necessarily go hand in hand. Thus, the powder’s flow requirements and attributes that enable sintering are in conflict. The powder that is ultimately used is a compromise between flow and sintering. There is no single grade (particle-size distribution) of powder used in RP/RM. Rather, those practicing the technologies make do with what is available. From our experience a good compromise entails a -270 or -325 mesh upper sieve cut. However, what appears to be more important is the amount of material that is <5 µm. This depends on the source of the powder, and an excess of fines will decrease the ability to doctor a precise and uniform layer. In this case air classification is used to remove the -5 µm fraction. This improves the flowability of the powder without negatively affecting its ability to sinter.
Q
What are the differences between thermal spray and cold spray and why is the powder for these technologies so expensive? Why is their use not more widespread? Ther mal spraying and cold spraying are processes that are used to coat substrates with a material by spraying the material, generally in a powder form, onto the substrate. Since powder is a necessary part of this technology one may consider this a PM technology and it should be embraced by the PM community. In thermal spraying the powder is heated to a molten or semi-molten state by injection into a heated plasma (plasma spraying) or a flame (flame spraying, such as oxy-fuel spraying). The molten droplets splat and bond to the substrate. Coatings can be built up to almost any thickness by building up multiple layers. Although the bonding between the coating and the substrate is mostly mechanical, the bonding between layers is metallurgical. In cold spraying the powder is not heated significantly. Rather, cold spraying equipment relies on accelerating the powder particles to high velocities in a gas stream (heating the gas increases the velocity). The kinetic energy of the particle is converted, in part, into thermal energy upon impact with the substrate and the particle splats and bonds to the substrate. The powder requirements for the two processes are different. For thermal spraying the particle size depends primarily on the heat transfer of the process (temperature, residence time in the hot zone, convection, etc.) and the thermal properties
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of the powder (conductivity, melting temperature). Particles that are too fine may evaporate in the thermal plume. Particles that are too coarse may not melt sufficiently to conform (splat) in the coating. Thus, the optimum particle-size distribution is tailored by cutting both sides of the as-atomized distribution, which effectively reduces the product yield. If there is a market for the under- and oversize powders then the cost of the thermal spray powder may be relatively inexpensive. However, many thermal-spray alloys are specific to thermalspray applications and there is no market for the off-size powders; thus, the cost will be high. In addition, the markets for these special alloys are not large in relation to press-and-sinter or MIM and do not benefit from economy of scale. Cold-spray powder suffers a similar fate. Most applications call for powder that is <25 µm to enable the particle’s velocity to be maximized in the gas stream. Again, this may represent a minor-
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ity of an as-atomized distribution and the price suffers the same lack of economy of scale. The markets for thermal-spray and cold-spray coatings are growing but there are a few stumbling blocks in the way. One of the most salient limitations to the use of these coating techniques is the lack of a sanctioned set of standards to guarantee minimum properties (bond strength and physical and chemical properties of the coating, etc.). There are numerous variations to each of the spraying technologies and the results are sensitive to the operating parameters, powder alloy and properties, equipment, substrate surface preparation and properties, etc. Most applications are qualified for a specific set of conditions and new uses must be qualified specifically for that particular application. ijpm Readers are invited to send in questions for future issues. Submit your questions to: Consultants’ Corner, APMI International, 105 College Road East, Princeton, NJ 08540-6692; Fax (609) 987-8523; E-mail:
[email protected]
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PM METALLOGRAPHY
CUTTING-EDGE PM METALLOGRAPHY Thomas F. Murphy, FAPMI*
Metallography is probably the most productive and efficient technique available to the materials scientist. This may sound like a bold statement, but in relation to physical, mechanical, and chemical testing the depth and breadth of the information obtained by metallography augments that derived from each of the three cited testing modes. For instance, in the powder metallurgy (PM) industry, where part properties are determined by the material density (level of porosity), chemical composition/alloying method, and microstructure, only metallography can provide information in all three areas. Results from the other testing regimes are extremely important and provide insight into specific performance characteristics of a composition and/or part design, but it is metallography that serves as the overarching entity that synthesizes the information. A complete picture of a PM part’s performance is not possible without input from all types of testing, but many key answers are obtained via metallography. In addition, metallography is the only discipline with the capability of characterizing and monitoring a PM part through its entire manufacturing history. Starting with examination of the loose powder raw material, through each manufacturing step, to the part being placed in service, and finally to possible failure, metallography can be used to characterize the changes within the microstructure, which explain the behavior of the part in service. Another important, but frequently overlooked, aspect of metallographic examination is the analysis of PM parts to determine why certain parts perform satisfactorily while others fail. This capability to analyze a wide variety of samples, from powder particles to sintered or powder-forged parts, sets metallography apart from the other test techniques. Thus, metallography is attractive to powder and parts makers, in addition to end users; it provides a common ground for discussion and understanding. Recent seminal advances have taken place in PM metallography and are covered in this Focus Issue. The authors invited to contribute, along with participants in the Special Interest Program (SIP) on “PM Metallography” at the PM2008 World Congress, have utilized a variety of sample types and analytical techniques to portray a compelling story of the utility and value of metallography, and its application to PM materials. Much of this information is at the forefront of currently available methods for analysis and
*Scientist, Research and Development, Hoeganaes Corporation, 1001 Taylors Lane, Cinnaminson, New Jersey 08077-2017; E-mail:
[email protected]
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sample preparation. One such presentation by Ver n Robertson, JEOL USA, demonstrated advancements in scanning electron microscopy (SEM) using low accelerating voltages. This approach provides a significant improvement in near -surface imaging and chemical analysis. When working with fine microstructural features and nanoscale distances, these techniques can be extremely valuable. This Focus Issue on PM Metallography illustrates a variety of applications in which metallography is used to characterize PM materials using several microscopy and sample preparation techniques. Carl Blais, Gilles L’Espérance, and Philippe Plamondon demonstrate the viability of transmission electron microscopy (TEM) to analyze powder particles and consolidated materials. Of particular interest is their use of a focused ion beam (FIB) to thin bulk specimens to electron transparency. PM materials are notoriously difficult to prepare for TEM examination and their detailed step-by-step description of the preparation procedure and subsequent imaging demonstrate an alternative to the tedious polishing of small samples that are transparent to the incident electron beam. Nik Chawla et al. present two techniques used to visualize and characterize the three-dimensional (3D) pore structure in PM materials. The first method uses serial sectioning involving consecutive grinding, polishing, and imaging of the same sample area to produce a series of images separated by the thickness of the material removed during preparation. The images are then placed into a stack using computer software and a model of the 3D microstructure is displayed. X-ray tomography, the second technique, is used to simulate the pore network. This method virtually eliminates serial-sectioning sample preparation by scanning the samples using high-voltage X-rays at predetermined intervals and creating a collection of the images of the individual slices. The set of images is again stacked and the 3D reconstruction of the pore network recreated. Both techniques provide high-resolution images of the interconnected pore network within a sintered
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compact. Data such as this are extremely valuable in understanding PM microstructures, in particular, the pore network/morphology. Paul Beiss and Stefan Lindlohr present a detailed statistical analysis of the pore structure and its relationship to fatigue strength in sintered iron bars. Using quantitative image analysis of metallographically prepared cross sections, they analyze the size and shape of the pore structure on the planar surfaces and use this information to predict improvements in fatigue strength through modification of the microstructure. Consideration is also given to the hardness of the material and how the combination of pore characteristics and hardness affect fatigue strength. The author presents a fractographic analysis of cracks and fracture surfaces. A combination of scanning electron microscopy (SEM) and light optical microscopy (LM) is used to quantify cracks contained within cross sections and exposed surfaces created by fracture. The amount of fracture on porous PM materials is estimated using SEM images on specimens pressed to several densities and sintered at different temperatures. In a second technique, crack profiles are examined to determine the amount of roughness resulting from fracture. In examining profiles, attention to detail is required to ensure that sample preparation of the fracture edge is accurate and small details are faithfully reproduced. A technique is also described to determine the location of cracks and their distribution within the microstructure. The purpose is to determine whether or not crack growth occurs preferentially within specific microstructural constituents. Using a combination of the three methods can help in determining how cracks form and materials eventually fail. In closing, let me reiterate that metallography is one of the most powerful tools we have to evaluate PM materials. In addition to the analytical capabilities cited, the aesthetic qualities of many microstructures are striking and pleasing to the eye. The natural beauty of well-prepared and etched or chemically stained sections is part of the artistry practiced by accomplished metallographers. ijpm
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THREE-DIMENSIONAL CHARACTERIZATION AND MODELING OF POROSITY IN PM STEELS Nikhilesh Chawla*, Jason J. Williams**, Xin Deng***, Casey McClimon****, Luke Hunter***** and S.H. Lau******
INTRODUCTION Sintered PM alloys normally contain residual porosity after sintering.1–7 The nature of the porosity is controlled by several processing variables such as green density, sintering temperature and time, alloying additions, and the particle size of the initial powders.1 The fraction, size, distribution, and morphology of the pores have a profound impact on mechanical behavior.2–7 In general, the porosity in sintered ferrous alloys is bimodal and can be classified as either primary or secondary. Primary porosity consists of larger pores, due mostly to powder packing characteristics, which result in less than complete densification during sintering. Secondary porosity consists of much smaller pores often caused by the transient liquid phase of one or more alloying additions that form during sintering. An example of an alloying element that forms a transient liquid phase in sintered steels is copper. Upon melting, the copper particles leave behind small, rounded “secondary” pores with a size approximating that of the original copper particles. The interconnectivity between pores is usually a function of the total amount of porosity. At porosity levels >5 v/o, the pores tend to be interconnected. Below this level the pores tend to be relatively isolated. Interconnected porosity causes an increase in strain localization at the relatively small sintered regions (necks) between particles, while isolated porosity results in a more homogeneous mode of deformation. It is also not uncommon for the pore distribution to be inhomogeneous. In this case, strain localization will take place at pore clusters.8 Thus, for a given level of porosity, interconnected pores are more detrimental than isolated pores and reduce macroscopic ductility to a greater extent. Characterization of the pore size and pore shape has routinely been carried out by sectioning and metallography, followed by image analysis and/or stereological approaches.9,10 Meaningful information can be obtained from these analyses, provided a representative section of the
Visualization of porosity in three dimensions (3D) is critical to a detailed understanding of the microstructure and properties of powder metallurgy (PM) steels. Two techniques are described to characterize porosity in 3D: serial sectioning and X-ray tomography. Both techniques can be used to visualize porosity, and to quantify pore fraction and the degree of pore interconnectivity. The use of image-based finite element simulations, using two-dimensional (2D) images of the microstructure, is also described. This is an invaluable tool in correlating pore structure with the mechanical properties of the steel.
*Professor, **Assistant Research Scientist, School of Materials, Fulton School of Engineering, Arizona State University, P. O. Box 876006, Tempe, Arizona 85287-8706; E-mail:
[email protected], ***Senior Engineer, Kennametal Inc., 205 North 13th Street, Rogers, Arkansas 72756 ****Engineer, Texas Instruments, Austin, Texas 78746, *****Engineer, ******Vice President, Xradia Inc., 5052 Commercial Circle, Concordia, California 94520
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microstructure is sampled to obtain meaningful statistics. This approach is conducted on 2D sections and images. Epoxy impregnation of the pores, followed by dissolution of the steel matrix, has also been used to visualize the nature of the porosity. 11–13 Two novel techniques have been exploited to visualize microstructures in 3D: (i) Serial Sectioning This technique includes serial mechanical polishing14–16 or focused ion beam (FIB) milling.17–19 2D images of the microstructure are obtained through the depth of the sample and fiducial marks, such as Vickers indentations (in the case of mechanical polishing), are used to determine the spacing between sections. Software is used to reconstruct a 3D virtual microstructure from the 2D images. The serial sectioning technique is, of necessity, destructive and is also time-consuming. (ii) X-ray Tomography20–22 In this technique X-rays are transmitted through the specimen at varying angles. Using inverse Fourier transforms, 2D “virtual slices” of the microstructure can be obtained. As in serial sectioning, software is used to reconstruct a 3D virtual microstructure from the 2D images. The technique is relatively simple and nondestructive, although materials with relatively high density are more problematic because of the high level of absorption of X-rays. Also, if the material consists of multiple phases of similar density, then there may not be enough contrast to differentiate between the phases. In such cases, other, more complex versions of the technique, such as refraction-based tomography 23 or holotomography 24 need to be employed. The X-ray tomography technique has recently been used to study evolution of the pore structure during the compaction and sintering of PM steels.21,22 Use of the 3D microstructures needs not be limited to visualization. The microstructures have also been used as a basis for sophisticated numerical models to simulate deformation behavior.25–28 In the current study, we show how serial sectioning and X-ray tomography can be used to visualize the 3D nature of porosity in PM steels. Quantitative analysis of the porosity, such as pore fraction and interconnectivity data, are also presented. The use of image-based finite element modeling to understand the influence of porosity on plastic deformation in these PM steels is also described.
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MATERIALS AND PROCESSING Powder mixtures of Fe-0.85 w/o Mo prealloy powder, 2 w/o Ni, and 0.6 w/o graphite were mixed and binder treated (Hoeganaes Corporation).29,30 The organic binder is used to bond the fine alloying particles to the larger iron particles, in order to decrease the diffusion distance between particles during sintering. The binder also improves the lubricity of the powder mixture during compaction prior to sintering. The average-size granule (prealloy powder, nickel, graphite, and binder) after the binder treatment was ~100–150 µm. Powders were compacted into rectangular blanks, and sintered at 1,120°C for 30 min in a 90 v/o N2-10 v/o H2 atmosphere. The binder was burned out during the initial ramp-up to the sintering temperature. Conventional compaction and sintering procedures were used to obtain a sintered density of approximately 7.0 g/cm3. Details of the processing parameters are cited elsewhere.31 SERIAL SECTIONING The following steps were used for the serial sectioning process, 3D visualization, and modeling, Figure 1. The PM samples were cut and mounted and a representative region of the microstructure seleted. Selection of the region of interest is important, but somewhat subjective. It is desirable to obtain a number of sections that encompass the porosity in the steel. An approximate volume of 110 × 140 × 15 µm3 was chosen, which yielded a total pore fraction similar to that calculated from the pore-free density of the PM steel. Fiducial marks, made by Vickers indentations, were used to measure the material thickness loss during polishing/grinding. Since the geometry of the indenter is known, the amount of material thickness removed can be calculated. The removal rate was taken as the average change in size of four indentations surrounding the region of interest, Figure 2. This removal rate was calibrated to give a thickness loss of about 0.6 µm per cycle. Quantifying material removal for serial sectioning in this manner has been used extensively due to its simplicity. The approximate depth (h) of a Vickers indentation can be calculated using the relationship: D h = ————— 2 tan (ϕ/2)
(1)
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Figure 1. Schematic of serial sectioning, 3D microstructure reconstruction, and visualization of porosity in PM steel
Figure 2. Optical micrograph showing porosity, Vickers hardness indentations used as fiducial marks, and region of interest in reconstructing 3D pore structure
where D is the average length of the indentation diagonals (D1 & D2) on a 2D projection and ϕ is the angle between the two diagonals. Cyclic polishing and imaging of the sample surface were then conducted to generate a series of Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
microstructural sections. The role of polishing in serial sectioning is important for two reasons: (a) to control the amount of material removed, and (b) to obtain a high-quality surface finish for microstructural characterization. In order to obtain a thickness loss ~0.6 µm per cycle, the sample was polished using 1 µm diamond paste on a nylon cloth. Polishing was conducted using an automatic polisher, with each step consisting of polishing for 15 min under a force of 5 N and at a speed of 25 rpm. After each polishing cycle, images of the microstructure were taken with an optical microscope. The sample was secured using a mounting fixture to minimize translational and rotational misalignments between sections. The microstructures were segmented into black and white images using conventional image analysis software (ImageJ, Bethesda, Maryland). The images were segmented and stacked in sequence and a 3D model was constructed using commercial reconstruction software (Mimics, Materialise, Ann Arbor, Michigan). Figure 3 shows a 3D reconstruction of the porosity, which is tortuous and highly interconnected. Figure 4 shows a comparison of the 2D porosity, taken at each of the 2D images prior to reconstruction, with the porosity measured from
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the 3D model. As expected, the 2D porosity, at any given section, underestimates or overestimates the actual 3D porosity. This is exacerbated by the highly irregular nature of the porosity in PM steels.
some disadvantages. Among them is that the process is laborious and destructive. X-ray tomography is a technique that is compatible with high-
X-RAY TOMOGRAPHY While serial sectioning is a powerful technique for generating virtual 3D microstructures, it has
Figure 3. 3D model reconstruction of porosity in PM steel using serial sectioning. Note the highly interconnected nature of the porosity
Figure 4. Comparison of 2D and 3D porosity measurements. The 2D measurements overestimate or underestimate the true v/o porosity
Figure 5. Schematic of X-ray tomography technique
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resolution 3D imaging and which eliminates cross sectioning, and allows for superior resolution, throughput, and image quality with minimal sample preparation. A schematic of the technique is shown in Figure 5. The PM steel sample was cylindrical (1.5 mm dia.), machined by electric discharge machining (EDM). The computed tomography image data were obtained using a MicroXCT -200 (Xradia, Inc., Concord, California). This system is a state-ofthe-art X-ray tomography tool capable of resolution down to 1.5 µm with high contrast. The samples were all scanned using the same conditions, namely, 150 kV and 4 W for X-ray source power; 20X objective providing a pixel size of 1.1 µm; 20 s per projection; 737 projections over 180°. A 100 µm steel filter was used to prefilter the X-ray beam to reduce beam-hardening arti-
Figure 6. (a) 2D slice obtained from X-ray tomography, and (b) accompanying segmented black-and-white image used in 3D reconstruction
Figure 7. Schematic of pore interconnectivity determined by a “region growth” algorithm. Shaded regions delineate growth during simulation to determine total interconnectivity
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
facts. The samples were scanned consecutively using Xradia’s recipe feature, which allows for multiple scans and reconstructions from a single set of input parameters. Samples were mounted on a highly stable and accurate stage, which is essential for high-resolution imaging. The projected views were reconstructed using Xradia's TXMReconstructor software and output as 16-bit Tiff format for further analysis and 3D reconstruction. About 1,000 slices were obtained, each spaced 1 µm apart. Figure 6(a) shows one such reconstructed 2D slice obtained from X-ray tomography. The slices were segmented into a black-and-white image, as shown in Figure 6(b). A 3D microstructure was obtained following a similar procedure to that used for the 2D slices obtained from serial sectioning. In addition to visualization of the porosity, the degree of interconnectivity of the pores was determined. Interconnectivity was determined by a region-growing algorithm. In this technique a gray level threshold is set for the pore, Figure 7, and a pixel (or voxel) is then selected within this pore. If the pixels connected to the selection are within the gray level threshold, the selection expands to include additional pixels. In 2D, square pixels are considered to be connected if pixels share an edge or a corner (8-connected). In 3D, cubic voxels are considered to be connected if voxels share a face, edge, or a corner (26-connected). Figures 8(a) and 8(b) show the 3D reconstructed volume of the porosity. The total pore volume is about 12 v/o, which correlates well based on a calculated sintered density of 7.0 g/cm3. The pore network in yellow consists of one single interconnected network, and makes up about 93% of the total pore volume. The green pores correspond to isolated pores. Two interconnected pores, shown in light blue and magenta, are highlighted in Figure 8(b). Note the highly tortuous and irregular nature of the porosity, which is clearly highlighted by the X-ray tomography technique, Figure 8(c). MICROSTRUCTURE-BASED SIMULATION OF DEFORMATION The virtual microstructure obtained by serial sectioning or X-ray tomography can also be used for microstructure-based modeling in either 2D or 3D. Figure 9 shows a meshed 2D image of the porosity in a PM steel used as the basis for finite element (FE) analysis of uniaxial loading. The analysis was two dimensional (2D) under plane-
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Figure 8. X-ray tomography 3D reconstruction of porosity in PM steel. Porosity in yellow shows one large interconnected network: (a) front view, (b) back view, (c) high-magnification image of two pore networks
strain conditions with the following boundary conditions: the left edge was fixed in the horizontal direction (U1 = 0) while load was applied to the right edge horizontally under displacement rate control. A 1% strain was applied to the model. A quadratic triangular mesh was employed in this simulation to conform to the irregular nature of the microstructure. A finer mesh was used in regions of pore clusters and a coarser mesh was employed in the matrix-rich areas. The constitutive behavior of the steel matrix was extrapolated from tensile test experiments. The microstructure of the steel matrix for different densities was relatively constant, based on our recent work,32,33 so we assume that the steel matrix is identical at all porosity levels. It should be noted that the 2D analysis effectively treats the pores as holes in the microstructure. Thus, this model is different from actual pores which are 3D and surrounded by matrix material. Nevertheless, the 2D analysis presented here shows the qualitative effects of the pore structure on localized plastic-strain initiation and its evolution around the pores. The macroscopic stress–strain behavior predicted by the model is shown in Figure 10. The stress–strain input for the pore-free PM steel (extrapolated from experimental stress–strain behavior at the three different sintered densities), is also shown. The simulations were conducted for steels at three different densities; thus, the microstructures are distinctly different. The PM steel with the lowest level of porosity exhibited a large number of small pores, compared with the PM steel with the highest porosity level. Note that even a slight decrease in porosity (4%–10%)
Figure 9. Mesh and boundary conditions used for FE analysis based on the actual microstructure in PM steel
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Figure10. Modeled stress–strain behavior by 2D FE analysis. A sharp decrease in strength is observed at 10 v/o porosity, consistent with the experimental data
results in a significant decrease in the strength of the PM steel, as was observed experimentally. This is rationalized by considering the equivalent plastic-strain evolution in each of the three microstructures shown in Figure 11. A significant amount of strain localization takes place in the sintered regions between the pores. In particular, networks of pores are effective in localizing the
strains in the steel ligaments between the pores. Thus, a small section of the microstructure is actually plastically deformed, so that a large portion of the material is essentially undeformed. The modeling results are confirmed by the experimental observations, namely, porosity leads to localized deformation and is inhomogeneous. 34–36 Strain intensification in the sintered ligaments between pores most likely serves as areas for crack initiation. With the onset of crack initiation, the large pores are linked, and the effective loadbearing area of the material will decrease locally and rapidly, resulting in fracture of the material. An increase in the level of porosity decreases the overall sintered ligament fraction and the spacing between pores, thus accelerating the intensification of strain in the matrix material. Our modeling also shows that plastic-strain intensification begins at the tips of the irregular pores in the microstructure. Vedula and Heckel34 compared the damage mechanisms between round and angular pores in materials with identical pore fractions and observed that highly localized slip bands formed at the sharp tips of angular pores, producing an uneven distribution of strain around angular pores. This resulted in highly localized and inhomogeneous plastic deformation compared with the deformation around round
Figure 11. Effective plastic strain contours in modeled microstructures as a function of sintered density: (a) 7.0 g/cm3, (b) 7.4 g/cm3, and (c) 7.5 g/cm3. Large interconnected pores cause strain intensification, while small, more homogeneously distributed pores contribute to more homogeneous deformation
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pores which was much more homogeneous. The distribution of the pores is also important since it has been shown that plastic deformation may initiate at pore clusters because of the higher localized stress intensity associated with these defects. The plastic-strain distribution in the modeled PM steel microstructure at a density of 7.5 g/cm 3 (Figure 11(c)) shows that when the pores are small and homogeneously distributed, the plastic-strain distribution is more uniform and the deformation is more uniformly distributed throughout the material. Most of the strain localization takes place at the shortest distance between pores or pore clusters. In particular, most of the plastic-deformation bands tend to be at an angle to the tensile direction. Thus, the orientation of pores with respect to the loading axis may also play a significant role in relation to plastic deformation. Thus, while the strength of the material is controlled by the fraction of pores, macroscopic ductility is also influenced by the size distribution, orientation, and degree of clustering of the pores, since the sintered ligaments in the PM steel control fracture of the material. An equally important result of the model is that, even in the highest-density material, a significant amount of strain intensification takes place at a single pore cluster in the microstructure, Figure 11(c). Thus, even when the overall amount of porosity is relatively low (4%–5%), strain intensification may take place around pore clusters. It follows that the homogeneity and distribution of the porosity is as important as the level of porosity in controlling the evolution of plastic strain, and thus the onset of crack initiation. SUMMARY Visualization of porosity in 3D is critical to a thorough understanding of the microstructure and properties of PM steels. Here we describe two techniques that can be used to visualize porosity in 3D: serial sectioning and X-ray tomography. Both techniques can be used to visualize porosity, as well as to quantify the pore fraction and degree of pore interconnectivity. The use of image-based finite element simulations, using 2D images of the microstructure, is also described. This is an invaluable tool in correlating the actual pore structure with the mechanical properties of the PM steel.
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ACKNOWLEDGMENTS The authors thank the Hoeganaes Corporation for providing the samples used in this study. Nikhilesh Chawla is indebted to Tom Murphy, Hoeganaes Corporation, for insightful discussions and for providing the references on epoxy impregnation. REFERENCES 1. A. Salak, Ferrous Powder Metallurgy, 1997, Cambridge International Science Publishing, Cambridge, UK. 2. A. Harboletz and B. Weiss, “Fatigue Behavior of Iron Based Sintered Material: a Review”, Int. Mater. Rev., 1997, vol. 42, pp. 1–44. 3. N. Chawla, S. Polasik, K.S. Narasimhan, M. Koopman and K.K. Chawla, “Fatigue Behavior of Binder-Treated P/M Steels”, Int. J. of Powder Metall., 2001, vol. 37, no. 3, pp. 49–57. 4. N. Chawla, T.F. Murphy, K.S. Narasimhan, M. Koopman and K.K. Chawla, “Axial Fatigue Behavior of Binder-treated versus Diffusion Alloyed Powder Metallurgy Steels”, Mater. Sci. Eng. A, 2001, vol. 308, no. 1–2, pp. 180–188. 5. N. Chawla, D. Babic, J.J. Williams, S.J. Polasik, M. Marucci and K.S. Narasimhan, “Effect of Ni and Cu Alloying Additions on the Tensile and Fatigue Behavior of Sintered Steels”, Advances in Powder Metallurgy and Particulate Materials—2002, compiled by V. Arnhold, C-L Chu, W.F. Jandeska Jr. and H.I. Sanderow, Metal Powder Industries Federation, Princeton, NJ, 2002, part 5, pp. 104–112. 6. S.J. Polasik, J.J. Williams and N. Chawla, “Fatigue Crack Initiation and Propagation in Binder -Treated Powder Metallurgy Steels”, Metall. Mater. Trans. A, 2002, vol. 33A, pp. 73–81. 7. N. Chawla, B. Jester and D.T. Vonk, “Bauschinger Effect in Porous Sintered Steels”, Mater. Sci. Eng. A, 2003, vol. 346, no. 1–2, pp. 266–272. 8. H. Danninger, G. Tang, B. Weiss and R. Stickler, “Microstructure and Mechanical Properties of Sintered iron Part II—Experimental Study”, Powder Metall. Int., 1993, vol. 25, no. 4, pp. 170–175. 9. A. Lawley and T.F. Murphy, “Metallography of Powder Metallurgy Materials”, Mater. Charac., 2003, vol. 51, no. 5, pp. 315–327. 10. “Particle Image Analysis”, ASM Handbook, Vol. 7: Powder Metal Technologies and Applications, ASM International, Materials Park, OH, 1998, pp. 259–273. 11. H. Danninger, D. Spoljaric, G. Jangg and B. Weiss, “Characterization of Pressed and Sintered Ferrous Materials by Quantitative Fractography”, Praktische Metallographie, 1994, vol. 31, no. 2, pp. 56–69. 12. H. Danninger, D. Spoljaric and B. Weiss, “Microstructural Features Limiting the Performance of PM Structural Parts”, Advances in Powder Metallurgy & Particulate Materials—1996, compiled by T. Cadle and K.S. Narasimhan, Metal Powder Industries Federation, Princeton, NJ, 1996, vol. 4, part 13, pp. 479–490. 13. T.F. Murphy and B. Lindsley, “Metallographic Analysis of PM Fracture Surfaces”, Advances in Powder Metallurgy and Particulate Materials—2007, compiled by J. Engquist
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14.
15.
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17.
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21.
22.
23.
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and T.F. Murphy, Metal Powder Industries Federation, Princeton, NJ, 2007, vol. 2, part 11, pp. 1–15. M.A. Dudek and N. Chawla, “Three-Dimensional (3D) Microstructure Visualization of LaSn3 Intermetallics in a Novel Sn-rich Rare-Earth Containing Solder”, Mater. Charac., 2008, vol. 59, no. 9, pp. 1,364–1,368. R.S. Sidhu and N. Chawla, “Three-Dimensional (3D) Microstructure Characterization of Ag3Sn Intermetallics in Sn-rich Solder by Serial Sectioning”, Mater. Charac., 2004, 52, pp. 225–230. N. Chawla, V.V. Ganesh and B. Wunsch, “ThreeDimensional (3D) Microstructure Visualization and Finite Element Modeling of the Mechanical Behavior of SiC Particle Reinforced Aluminum Composites”, Scripta Mater., 2004, vol. 51, no. 2, pp. 161–165. A.J. Kubis, G.J. Shiflet and R. Hull, “Focused Ion-Beam Tomography”, Metall. Mater. Trans., 2004, vol. 35, no. 7, pp. 1,935–1,943. B.J. Inkson, T. Steer, G. Möbus and T. Wagner, “Subsurface Nanoindentation Deformation of Cu-Al Multilayers Mapped in 3D by Focused Ion Beam Microscopy”, Journal of Microscopy, 2001, vol. 201, no. 2, pp. 256–269. C. Holzapfel, W. Schäf , M. Marx , H. Vehoff and F. Mücklich, “Interaction of Cracks With Precipitates and Grain Boundaries: Understanding Crack Growth Mechanisms Through Focused Ion Beam Tomography”, Scripta Mater., 2007, vol. 56, no. 8, pp. 697–700. F. Beckmann, R. Grupp, A. Haibel, M. Huppmann, M. Nöthe, A. Pyzalla, W. Reimers, A. Schreyer and R. Zettler, “In-Situ Synchrotron X-ray Microtomography Studies of Microstructure and Damage Evolution in Engineering Materials”, Adv. Eng. Mater., 2007, vol. 11, pp. 939–950. A. Vagnon, O. Lame, D. Bouvard, M. Di Michiel, D. Bellet and G. Kapelski, “Deformation of Steel Powder Compacts During Sintering: Correlation Between Macroscopic Measurement and in situ Microtomography Analysis”, Acta Mater., 2006, vol. 54, no. 2, pp. 513–522. S. Giménez, A. Vagnon, D. Bouvard and O. Van der Biest, “Influence of the Green Density on the Dewaxing Behaviour of Uniaxially Pressed Powder Compacts”, Mater. Sci. Eng., 2006, vol. 430, no. 1–2, pp. 277–284. M.P. Hentschel, A. Lange, B.R. Muller, J. Schors and K.W. Harbich, “X-ray Refraction Computer -Tomography”, Materialprufung, 2000, vol. 42, pp. 217–221. P. Kenesei, H. Biermann and A. Borbely, “StructureProperty Relationship in Particle Reinforced Metal-Matrix Composites Based on Holotomography”, Scripta Mater., 2005, vol. 53, no. 7, pp. 787–791. N. Chawla and K.K. Chawla, “Microstructure-Based Modeling of Deformation in Particle Reinforced Metal Matrix Composites”, J. Mater. Sci. – 40th Anniversary
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Issue (1966–2006), 2006, 41, pp. 913–925. 26. N. Chawla and R.S. Sidhu, “Microstructure-Based Modeling of Deformation in Pb-free Solders”, J. Mater. Sci—Materials in Electronics, Special Issue on Pb-free Solders, 2007, vol. 18, pp. 175–189. 27. N. Chawla, R.S. Sidhu and V.V. Ganesh, “Three Dimensional (3D) Visualization and Microstructure-Based Finite Element Modeling of Particle Reinforced Composites”, Acta Mater., 2006, vol. 54, pp. 1,541–1,548. 28. N. Chawla and X. Deng, “Microstructure and Mechanical Behavior of Porous Sintered Steels”, Mater. Sci. Eng., 2005, vol. A390, pp. 98–112. 29. F.J. Semel, “Properties of Parts Made from an Ancorbond® Processed Carbon–Nickel–Steel Powder Mix (FN-0208)”, Advances in Powder Metallurgy & Particulate Materials—1989, compiled by T.G. Gasbarre and W.F. Jandeska, Metal Powder Industries Federation, Princeton, NJ, 1989, vol. 1, pp. 9–23. 30. S.H. Luk and J.A. Hamill Jr., “Dust and Segregation-Free Powders for Flexible P/M Processing”, Advances in Powder Metallurgy & Particulate Materials—1993, compiled by A. Lawley and A. Swanson, Metal Powder Industries Federation, Princeton, NJ, 1993, vol. 1, pp. 153–168. 31. G. Poszmik and S.H. Luk, “Binder Treated Products for Higher Densities and Better Precision”, Advances in Powder Metallurgy & Particulate Materials—2003, compiled by R. Lawcock and M. Wright, Metal Powder Industries Federation, Princeton, NJ, 2003, part 3, pp. 33–44. 32. X. Deng, G.B. Piotrowski, N. Chawla and K.S. Narasimhan, “Fatigue Crack Growth Behavior of Hybrid and Prealloyed Sintered Steels, Part I: Microstructure Characterization”, Mater. Sci. Eng., 2008, vol. A491, pp. 19–27. 33. X. Deng, G.B. Piotrowski, N. Chawla and K.S. Narasimhan, “Fatigue Crack Growth Behavior of Hybrid and Prealloyed Sintered Steels, Part II: Fatigue Behavior”, Mater. Sci. Eng., 2008, vol. A491, pp. 28–38. 34. K.M. Vedula and R.W. Heckel, “Structural-Property Relations for the Tensile Behavior of Single Phase Ductile Sintered Materials”, Modern Developments in Powder Metall., edited by H. Hausner, H.W. Antes and G.D. Smith, Metal Powder Industries Federation, Princeton, NJ, 1980, pp. 759–777. 35. W.A. Spitzig, R.E. Smelser and O. Richmond, “The Evolution of Damage and Fracture in Iron Compacts with Various Initial Porosities,” Acta Metall., 1988, vol. 36, p. 1,201. 36. G. Straffelini and A. Molinari, “Evolution of Tensile Damage in Porous Iron,” Mater. Sci. Eng., 2002, vol. 334, no. 1–2, pp. 96–103. ijpm
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PM METALLOGRAPHY
CHARACTERIZATION OF POWDERS AND PM COMPONENTS UTILIZING TRANSMISSION ELECTRON MICROSCOPY Carl Blais,* Gilles L’Espérance, FAPMI,** and Philippe Plamondon***
INTRODUCTION The properties of a material are determined primarily by its microstructure. Therefore, quantitative microstructural characterization is important in the improvement of existing materials and the development of new ones. Property enhancement is now performed at the microscale level and even at the nanoscale level. Concentration gradients in diffusion-alloyed powders,1 precipitation hardening of PM aluminum,2 and second-phase particles used to improve wear resistance3 are examples existing on a size scale beyond the capabilities of conventional characterization methods. Therefore, in order to fully characterize such fine features, it is necessary to use a technique that yields a high spatial resolution, namely, TEM. Not only does TEM offer a spatial resolution that can resolve atoms, it also permits chemical characterization to be performed using energy-dispersive X-ray spectrometry (EDS) and electron energy-loss spectroscopy (EELS).4 TEM also offers the ability to acquire electron-diffraction patterns that provide information about the crystalline structure of the region characterized. Therefore, by combining the information available from imaging, EDS, EELS, and electron diffraction, it is possible to completely characterize nanometric microstructural features.5 Although TEM characterization offers these advantages, specimen preparation is challenging. Indeed, suitable TEM specimens have to be thin enough to be “transparent” to the incident electron beam. This means that the region of interest in the specimen has to be ~200 nm thick or less. In order to thin specimens for TEM characterization, several techniques are available. Among the most popular are ion milling,6 electrochemical thinning,6 and ultramicrotomy.7 Unfortunately, these techniques are not well suited for the preparation of PM specimens since they require mechanical thinning prior to the final step in thickness reduction. This involves holding in place a small specimen (3 mm in dia.) while it is being polished using abrasive media such as SiC-coated films or diamond sus-
Characterization of the fine microstructure of materials by transmission electron microscopy (TEM) is generally perceived as a long and tedious task. It is particularly true for powder metallurgy (PM) components in which the presence of porosity makes sample preparation even more difficult. Examples are presented in which the focused ion beam technique is used to prepare PM specimens for characterization by TEM. These examples illustrate the benefits of this technique and TEM in terms of sample quality and the added information that it yields. A detailed characterization of secondphase particles in resulfurized (prealloyed) PM components is presented.
*Professor, Department of Mining, Metallurgical and Materials Engineering, Université Laval, Québec City, QC, Canada; E-mail:
[email protected], **Professor, ***Research Assistant, Center for Microscopy and Characterization of Materials-(CM)2–École Polytechnique de Montréal, Montréal, QC, Canada
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pensions. The presence of porosity makes it difficult to obtain the desired wedge-shaped thin areas. Indeed, PM specimens have a tendency to collapse due to their poor structural integrity caused by the presence of pores. Second-phase particles have always played an important role in the overall performances of materials. Over the last two decades, we have seen the emergence of a new field of research in metallurgy called “inclusion engineering”8 which topic overlaps chemical and physical metallurgy since, in the case of steels, the liquid metal is precisely controlled to obtain specific second-phase particles that improve the properties of the end product. Examples of “inclusion engineering” can embrace microalloyed steels,9 welding,10 and machining.11 PM is not indifferent to “inclusion engineering.” Indeed, second-phase particles are often added to PM systems to improve their machining response. Thus, several compounds such as MnS,12 BN,13 MgSiO3,14 and free graphite15 have been used as additives in the form of second-phase particles to lower the forces necessary to cut the material during machining operations. Another example in which secondphase particles play an important role in improving the mechanical performances of PM components is the precipitation of carbides (precipitation hardening) in low-alloy and/or sinter-hardenable materials.16 Similarly, carbides or nitrides which are stable at high temperature can be introduced in PM components, either by premixing (carbide and steel powders) or prealloying to increase the wear resistance of the finished part.3 In the latter example, control of the chemistry of the precipitates is extremely important as it dictates the solubility of the compounds at high temperature (e.g., in austenite) and, therefore, the overall wear performance of the component. These examples all have in common the fact that the size of the second-phase particles is smaller than a few micrometers in diameter. Thus, the characterization of such small features is clearly beyond the capability of optical microscopy. Even scanning electron microscopy (SEM) with EDS can only be used to qualitatively characterize the chemistry of the second-phase particles due to their small size.17 Therefore, TEM is currently the only technique that yields a sufficiently high spatial resolution to fully characterize the size, shape, chemistry, and crystal structure of such small features. Sample preparation in TEM is the key to successful characterization. It
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has also been shown that the focused ion beam technique (FIB) is the sample preparation method that is best suited for TEM work on PM powders and components. Here we highlight the benefits of the FIB technique as applied to the characterization of second-phase particles in resulfurized (prealloyed) powders and components. METHODOLOGY Focused Ion Beam Sample Preparation Figure 1 is a schematic diagram of a focused ion beam (FIB) system that can be used to prepare thin foils for TEM characterization. The source of ions consists of a tungsten needle that is wetted by liquid gallium. An electric field is applied to the tip of the needle by the anode to extract metallic ions (Ga+), which are then accelerated by applying a positive potential to the accelerating anode. The selection of the accelerating voltage determines the kinetic energy supplied to the metallic ions. The extracted ions are then collimated and focused on the sample using electrostatic lenses.
Figure 1. Schematic representation of the main components of a FIB system
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Finally, scanning coils are used to raster the ion beam on the surface of the sample. As the ions interact with the sample, they transfer part of their kinetic energy to the atoms located in the interaction volume. These atoms are then sputtered in the form of ions or electrically neutral atoms. Thus, by precisely controlling the position of the beam, it is possible to select a specific area of interest and to mill this small area to obtain a specimen sufficiently thin for TEM characterization.
Figures 2 through 10 illustrate the primary steps involved in preparing a thin foil for TEM characterization using the FIB technique. Figure 2 is a low-magnification scanning electron micrograph (SEM) secondary electron image (SEI) of particles of a resulfurized steep powder, prealloyed with MnS particles. The objective of the characterization work was to investigate the chemistry of the MnS particles and their possible epitaxial relationships with the iron matrix after
Figure 2. Steel particles from which a thin foil will be extracted. SEM/SEI
Figure 3.Tungsten layer deposited on top of the surface of interest. SEM/SEI
(a)
(b)
Figure 4. Milling of a staircase in the vicinity of the area of interest (arrowed). (a) top view, (b) plan view. SEM/SEI
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atomization and prior to sintering. The arrow in Figure 2 delineates an area of interest that appears suitable for specimen preparation for TEM characterization. The first step in the procedure involves the deposition of a tungsten layer on top of the area of interest, Figure 3. The heavy metal deposit serves to protect the area during subsequent ion milling operations. Once tungsten is deposited, the ion milling operation can proceed. Starting ~30 µm from the area of interest, the iron matrix is sputtered by the ion beam. The sputtering depth of the beam is increased, by changing the accelerating voltage of the system, as it gets closer to the area of interest. This results in a hole with the shape of a staircase milled from the surface of the sample to a depth approximately 10 µm near one of the edges of the future specimen, Figure 4. Next, the contours of the specimen are milled, Figure 5, taking care to leave a bridge between the bulk of the sample and the specimen to prevent the latter from falling once the bottom is milled from the rest of the particle, Figure 6. Prior to milling the bridge that constitutes the last link between the thin foil in preparation and the bulk of the particle, a probe is attached to the specimen. The solid contact between the probe and the specimen is obtained by depositing a bead of tungsten that acts as a weldment, Figures 7 and 8. The bridge is then milled to completely detach the specimen from the rest of the powder particle. The probe is then used to
move the specimen and position it on a TEM grid designed to hold the specimen for the final preparation stage, Figure 8. A schematic representation of the grid used is given in Figure 9. One important feature of the FIB technique is that the TEM grid that is used for the final thinning operations is located in a TEM specimen holder. Thus, once the specimen is sufficiently thin, it can readily be introduced in the TEM without further manipulations. This is an advantage of the technique since monitoring of the specimen thickness can be performed without having to touch the specimen with tweezers. Once the specimen is attached to the TEM grid, the probe is cut from the specimen and the final ion milling operations are carried out. Finally, when the specimen is judged acceptable, Figure 10, it is introduced in the TEM and the characterization work can begin. Figure 11 is a low-magnification TEM micrograph of the final result. This shows that the useful area of analysis is large and the thickness profile is relatively constant throughout the surface of the specimen.
Figure 5. Milling of contours on the specimen. A bridge is left to support the specimen once the bottom is milled. SEM/SEI
Figure 6. Milling of bottom area of the specimen. At this point, the specimen is held in place only by the bridge (Figure 8). SEM/SEI
TEM Characterization An analytical transmission electron microscope* was used to characterize the specimens prepared using FIB. The TEM is a Phillips CM30 (maximum accelerating voltage of 300 kV), with a *Center for Microscopy and Characterization of Materials, (CM)2, École Polytechnique de Montréal
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Figure 7. Fixing of probe used to hold and move the specimen once it is completely detached from the rest of the particle. SEM/SEI
Figure 8. Placement of specimen on TEM grid using the probe. Tungsten weldments are used to hold the specimen in place before releasing the probe by cutting its tip. SEM/SEI
Figure 9. Schematic representation of TEM grid used in final ion thinning operation and characterization in the TEM
lanthanum hexaboride (LaB6) emitter. The TEM was also equipped with an EDS spectrometer (Link AN-10000) and a parallel EELS system (Gatan 666). As an example, an FC-0205 + 0.5 w/o prealloyed MnS powder mix was used to prepare cylindrical specimens (height 50.8 mm (2.0 in.) × dia. 38.1 mm (1.5 in.)) pressed to a green density of 6.8 g/cm 3 . The specimens were sintered at 1,121°C (2,050°F) for 25 min (at temperature) in an atmosphere of 90 v/o N2/10 v/o H2. One of the specimens was then placed in a CNC-lathe and turned by orthogonal cutting. Chips formed during the process were collected and prepared for TEM characterization using the FIB sample preparation technique. Thin foils having a thickness of 50 to 80 nm were obtained. Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
Figure 10. Final appearance of specimen ready for characterization in the TEM. SEM/SEI
EXAMPLES MnS Particles in Prealloyed Resulfurized Powder The goal of this work was to identify the micromechanisms operative during chip formation of components made from prealloyed resulfurized powders. The most widely accepted mechanisms advanced to explain the role of MnS particles as machining aids can be summarized as follows: (1) stress concentrators in the primary shear
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Figure 11. Low-magnification micrograph of a single powder particle of resulfurized steel prepared using the FIB technique. TEM
zone,18,19 and (2) lubricant at the tool surface facilitating chip movement and preventing seizure.20,21 It is also recognized that only 20% of the total energy in cutting a material is spent at the chip/tool interface,22 the majority being consumed in the primary shear zone ahead of the tool tip. Two types of MnS particles were observed in the thin foils. The first type, shown in Figure 12, consists of relatively spherical particles. These particles were generally accompanied by the presence of voids formed at both ends of a plane oriented along the shear-strain axis. Moreover, this type of particle was normally not pure MnS. Based on the X-ray spectra in Figures 12(b) and 12(c), the particles were primarily dual phase, namely MnS and MnSiO3. The image and spectra in Figure 12 were acquired by scanning transmission electron microscopy (STEM). An example of the second type of MnS particles is shown in Figure 13. The montage of micrographs reveals that this latter type of particle is elongated. X-ray spectra acquired at several locations along the length of
Figure 12. (a) dual-phase particles of MnS/MnSiO3, (b) representative X-ray spectrum from areas delineated “A”, (c) representative X-ray spectrum from areas delineated “B”, “C” delineates crack initiation
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Figure 13. (a) Mosaic of TEM micrographs showing the second type (elongated) of MnS particles, (b) representative X-ray spectrum from elongated MnS particles
these particles confirms that they are singlephase MnS (Figure 13(b)). It has been shown that the effectiveness of MnS particles in reducing the forces necessary to promote chip formation in the primary shear zone improves when the particles are dual phase, (MnS + MnSiO3) instead of pure MnS.20 The reason is that the presence of a harder MnSiO 3 phase reduces the plasticity of the particles, as illustrated in Figure 12(a). Therefore, the dual-phase particles are more effective stress concentrators than pure MnS in the primary shear zone, leading to a reduction of the forces required to develop microcracks and initiate chip formation. Although the pure MnS particles are not as effective stress concentrators as their dual-phase counterparts, they do act as thermoplastic shear bands that concentrate plastic flow in small volumes within the chips. As seen in Figure 14, the flow of material during chip formation is far from homogeneous. Indeed, although most of the deformation takes Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
place along the primary shear plane, considerable deformation also takes place in the secondary shear zone. The role of the dual-phase particles (MnS + MnSiO3) in the secondary shear zone is not as significant as it was in the primary shear zone. This is due to the fact that their work as stress concentrators is mostly done once they move from the primary to the secondary shear zone. Furthermore, the movement of the chip on the tool surface causes a change in the stress distribution. Indeed, the plane on which the shear stress was applied in the primary shear zone tends to be partially in compression in the secondary shear zone (zone A, Figure 3). Proof of the change in stress distribution from shear to compression is given by the occasional presence of built-up edges formed on the surface of the cutting tool (not present in this case). Thus, the voids that were formed in the primary shear zone, and that did not grow to produce a significant crack, are forced to close due the compressive stress
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Figure 15. Representative micrograph from secondary shear zone (chip/tool interface) showing grain structure. TEM
Figure 14. Cross section of representative chip. Region “A” delineates the flow zone. Arrow “B” delineates the secondary shear zone (chip/tool interface). Optical micrograph; etchant nital 2%
generated in the secondary shear zone. On the other hand, the role of the pure MnS particles in the secondary shear zone is important. Their propensity to undergo extensive deformation tends to minimize seizure3 of the flow zone as it moves on the tool’s surface (Figures 13 and 14). This behavior of the material in the flow zone increases the bulk speed of the chip in relation to the tool while decreasing the sliding forces applied on the tool. Figures 15 and 16 are TEM micrographs and electron diffraction patterns of a section of the secondary shear zone (chip/tool interface) of a chip. It can be seen that the grain size of the material in this region is significantly smaller than that of the original sintered material. The TEM micrographs in Figure 16 indicate that the material in the secondary shear zone is susceptible to dynamic recrystallization, due to severe deformation and a temperature rise. Indeed, the difference in orientation across the grain boundaries of neighboring grains is several degrees, indicating that high-angle grain boundaries exist, typical of recrystallization. One exception is seen from a comparison of the electron diffraction patterns of grain 1 and 6. Here the dif-
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ference in orientation between these two grains is small, indicating that grain 6 is a sub-grain. The results presented above show that prevention of seizure in the flow zone by the presence of pre-alloyed MnS particles is crucial to good machining performances since the softening induced by the elevated temperature could well be counterbalanced by the increase in yield strength induced by reduction of the grain size. In that case, both the higher temperature and the larger forces induced by the small recrystallized grains would accelerate tool wear. Finally, an interesting observation from the results presented above is that the mechanism of tool lubrication attributed to the presence of MnS particles may not play such an important role as it was initially thought. Indeed, since a flow zone is always present during chip formation,21 the role of the MnS is more to reduce the width of that region and therefore the risk of chip seizure in addition to lubricating the tool surface. CONCLUSIONS • Sample preparation using the FIB technique is well suited to the study of PM powders and components by TEM. The main advantages of the technique are: the ability to precisely select the specimen to be characterized from a small predetermined volume, and to achieve thickness homogeneity in the final sample. Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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Figure 16. Representative TEM micrograph and electron-diffraction patterns from secondary shear zone (chip/tool interface) showing significant crystallographic disorientation between neighboring grains
• In the primary shear zone, dual-phase inclusions (MnS + (MnSiO3)) are more effective stress concentrators than pure MnS particles. • In the secondary shear zone, pure MnS particles play an important role as they reduce seizure in the flow zone at the surface of the cutting tool, thereby minimizing sliding forces. • Lubrication of the cutting tool by MnS does not appear to be the main reason for improvement in machinability induced by the presence of prealloyed MnS particles. REFERENCES 1. G. L’Espérance, S. Martel and A. de Rege, “Detailed Microstructure Characterization of Recently Developed Partially Alloyed Powders”, Advances in Powder Metallurgy and Particulate Materials—1996, compiled by T.M. Cadle and K.S. Narasimhan, Metal Powder Industries Federation, Princeton, NJ, 1996, vol. 1, part 1, pp. 231–249. 2. W.H. Hunt, “New Directions in Al-Based P/M Materials for Automotive Applications”, Powdered Metal Applications, SP-1535, Soc. Automotive Engineers, Warrendale, PA,
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
2000, pp. 1–7. 3. C. Blais and G. L’Espérance, “The ‘MicrostructureMechanical Properties-Machinability’ Relationships of Sinter-Hardening Powders”, Special Interest Program: Advances in P/M Machining, 2001, PM2TEC2001, Metal Powder Industries Federation, Princeton, NJ, oral presentation. 4. R.F. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope, 1989, Plenum Press, New York, NY. 5. C. Blais, G. L’Espérance and É. Baril, “Characterization of 25-75 nm Phases in Inclusions found in C-Mn Steel Weldments Containing Titanium”, J. Microsc, 1989, vol. 189, part 3, pp. 249–262. 6. D.B. Williams and C.B. Carter, Transmission Electron Microscopy—A Textbook for Materials Science, 1996, Plenum Press, New York, NY. 7. G. McMahon and T. Mallis, Microsc. Res. Tech., 1995, 31, pp. 267–274. 8. C. Blais and G. L’Espérance, “Characterization of Inclusions in Steel, When Not Only Size Matters”, CIM Bulletin, Canadian Institute of Mining, Metallurgy and Petroleum, Montréal, PQ, Canada, 1998, vol. 91, no. 1,021, pp. 116–120. 9. T. Gladman, The Physical Metallurgy of Microalloyed Steels, 1997, Maney, London, UK
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10. G.M. Evans, “Microstructure and Properties of Ferritic Steel Welds Containing Ti and B”, Weld. J., 1996, (8), 251s–260s. 11. S.V. Subramanian, H.O. Gekonde, X. Zhang and J. Gao, “Inclusion Engineering of Steels for High Speed Machining”, CIM Bulletin, Canadian Institute of Mining, Metallurgy and Petroleum, Montréal, PQ, Canada, 1998, vol. 91, no. 1,021, pp. 107–115. 12. L. Roy, G. L’Espérance and P. Lambert, “Prealloyed Mn/S Powders for Improved Machinability in PM Parts”, Met. Powder Rep., 1989, vol. 44, no. 2, pp. 116–119. 13. F. Chagnon, Powder Metallurgy Machinability Seminar, Metal Powder Industries Federation, Baltimore, MD, August 1996. 14. K. Hayashi, H. Shikata, Y. Ikenoue, K. Ishii, K. Chikahata and G. Goyo, “Enhanced Machinability of Valve Guides Made From P/M Materials”, Advances in Powder Metallurgy and Particulate Materials—1996, compiled by T.M. Cadle and K.S. Narasimhan, Metal Powder Industries Federation, Princeton, NJ, 1996, vol. 4, part 13, pp. 117–121. 15. S. Uenosono, S. Unami and K. Ogura, “A New Improvement in the Machinability of P/M Steel Due to Retained Graphite Particles”, Advances in Powder Metallurgy and Particulate Materials—1995, compiled by M. Phillips and J. Porter, Metal Powder Industries Federation, Princeton, NJ, 1995, vol. 2, part 8, pp. 171–176. 16. G. L’Espérance, “Design and Optimization of Sinter Hardening Processes”, Seminar on Sinter -Hardening, Metal Powder Industries Federation, April 13–14, Cleveland, OH, 1999. 17. C. Blais, G. L’Espérance and A. de Rege, “State of the Art Characterization of Second Phase Particles Found in P/M Powders and Parts”, Advances in Powder Metallurgy and Particulate Materials—1997, compiled by R.A. McKotch and R. Webb, Metal Powder Industries Federation, Princeton, NJ, 1997, vol. 2, part 15, pp. 3–15. 18. C. Blais, G. L’Espérance and I. Bourgeois, “Characterisation of Machinability of Sintered Steels During Drilling Operations”, Powder Metal., 2001, vol. 44, no. 1, pp. 67–76. 19. R. German, Powder Metallurgy Science, Second Edition, Metal Powder Industries Federation, Princeton, NJ, 1994, p. 346. 20. M.C. Shaw, Metal Cutting Principles, Oxford University Press, Oxford, UK, 1989. 21. P. Plamondon, G. L’Espérance and C. Blais, “Optimization of Pre-alloyed MnS Steel Powders for Improved Machinability”, Advances in Powder Metallurgy and Particulate Materials—2001, compiled by W.B. Eisen, and S. Kassam, Metal Powder Industries Federation, Princeton, NJ, 2001, part 6, pp. 29–39. 22. E.M. Trent and P.K. Wright, Metal Cutting, Butterworth Heinemann, Woburn, MA, 2000. ijpm
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PM METALLOGRAPHY
POROSITY STATISTICS AND FATIGUE STRENGTH OF SINTERED IRON Paul Beiss, FAPMI,* and Stefan Lindlohr**
INTRODUCTION For many years observations have been reported which suggest a relationship between large pores and the fatigue strength of sintered iron and steel. Plane-bending-endurance limits were measured with sponge iron and a water-atomized iron powders that had been compacted and sintered to a range of densities.3,4 When fatigue strength was plotted versus sintered density, the sponge iron material exhibited a higher performance compared with the water-atomized material at the same density. This effect was attributed to the smaller and more numerous pores in the sponge iron. The pore-size difference was documented by surface-roughness measurements and interpreted by utilizing fracture mechanics. When the endurance limits of both materials were plotted versus the inverse of the square root of the surface roughness Rt, a common straight line with minor scatter was obtained. Cimino et al.5 performed rotating-bending-fatigue tests on a Fe-CuC steel. The copper was added to the mix in three different size fractions. Irrespective of density or sintering conditions, the material with the coarsest copper powder and the largest secondary pores always exhibited the lowest fatigue strength. Using quantitative image analysis of large numbers of pores Beiss and Dalgic6 established a relationship between pore size and the irregularity of the pores for several carbon-free irons and steels. Statistically, the largest pores also exhibited the highest degree of irregularity which was expressed by a shape (or form) factor F given by 4πA/P2, where A is the pore area and P is the pore perimeter in a polished unetched section, Figure 1. The evaluation was restricted to the most irregular pores (2%) which were also the largest pores or delaminations from ejection. A good correlation was found between the planebending-fatigue strength and 4πA/P 2 for the pores selected. Murakami1 has shown by fracture mechanics that, in pore-free materials, there is a proportionality between the fatigue strength of a set of specimens and the material hardness divided by the sixth root of the largest defect size within the highly loaded volume of the specimens. To estimate the largest defect, Murakami1 applied Gumbel’s statistics of extremes 2 by analyzing a multitude of metallographic microsections from which only the largest defect per field was selected.
Murakami1 has suggested that Gumbel’s statistics of extremes2 can be applied to structural defects and inclusions. We attempt to interpret Murakami’s ideas for porous sintered steels by considering large pores as structural defects. Evaluating the largest pores quantitatively should yield a correlation between fatigue strength and pore dimensions. To test this hypothesis, plane bending-fatigue specimens (ISO 3928) were pressed from nominally pure iron and sintered at 1,120°C and 1,280°C. After fatigue testing, the specimens were evaluated by quantitative image analysis of the porosity in the waist section of the samples where fracture occurred. From each image the largest pore or oxide was selected and subjected to Gumbel’s distribution which gives a characteristic largest pore size λ and a distribution parameter δ for each material. λ was assumed to be representative of the material condition. The fully reversed bending endurance limit was found to be proportional to the Brinell hardness divided by the sixth root of λ. Thus, only a few extremely large pores out of thousands of pores appear to be important in affecting the fatigue behavior of sintered materials. If this relationship between fatigue strength, macroscopic hardness, and extreme pore size can be generalized, doubling the hardness should double the fatigue strength. Reducing the size of the largest pores by a factor of two would increase the endurance limit by only 11%.
*Professor, **Metallography Staff, Institute for Materials Applications in Mechanical Engineering, RWTH Aachen University, Augustinerbach4, D-52062 Aachen, Germany; E-mail:
[email protected]
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Figure 1. Primary pore characteristics in quantitative image analysis
The small number of largest defects from a representative number of fields characterizes the defect structure of the material and can be used to estimate the largest defect for deviating highly loaded volumes. In addition, Murakami assumed a proportionality between fatigue strength and the Vickers hardness of the pore-free material up to 400 HV. In this way, he was able to estimate the lower limit of fatigue strength of a variety of steels and nonferrous alloys with reasonable accuracy. The present study is an attempt to apply this general idea to the porosity in sintered iron. EXPERIMENTAL PROCEDURE Three types of pure iron were chosen for this investigation: a round, double-reduced wateratomized powder (ABC 100.30), a pure broken sponge of high irregularity (MH 65.17), and a sponge iron powder of intermediate irregularity (NC 100.24). All three grades were supplied by Höganäs AB, Sweden. The powders were mixed with 1 w/o microwax and compacted in the form of unnotched plane-bending-fatigue test specimens (ISO 3928). Different compaction pressures were utilized to achieve a range of green densities. Half the specimens were sintered at 1,120°C and the other half at 1,280°C in an industrial furnace, mostly in an atmosphere of 70 v/o N2/30 v/o H2. The NC 100.24 specimens were sintered in error at 1,120°C in 95 v/o N2/5 v/o H2 plus 100 L/h CH4. These test series are identified in Table I. The specimens were tested in fully reversed plane bending up to 107 cycles. Generally 60 samples were evenly distributed at six stress levels to achieve three levels with fractures and runouts
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for the statistical evaluations. Since die filling results in a lower filling density, and subsequently lower as-compacted and sintered densities in thin cross sections,7–10 the densities were determined from small sections of about 2 g to 3 g cut from the waist of the hourglass specimens. Ten samples each were taken, the results of which were averaged. The hardness was measured in the waist zone for the same reasons. The Brinell hardness method was used with a 2.5 mm dia. cemented carbide ball and loads of 31.25 kgf or 62.5 kgf, depending on the hardness reading. Because of the density distribution along the specimen axis the metallographic investigation was restricted to the thinnest cross section of the runout specimens. In each material condition five samples were taken, resin mounted, impregnated with epoxy, ground and prepolished, and vacuum impregnated with epoxy for a second time before final grinding and polishing. From each sample, ten images were taken at a magnification of 100:1. This gave fifty fields for each material condition which corresponded to an analyzed area of 67.62 mm2 with a resolution 1,536 × 1,160 pixels. Pores smaller than five pixels or below 3.794 µm2 were excluded. The total number of pores was between ~100,000 for the low-density wateratomized powder and 440,000 for the higher-density sponge iron after sintering at 1,120°C. The higher sintering temperature reduced the number of pores to 90,000 and 360,000, respectively. From each of the fifty fields for each S-N curve the largest pore was selected for analysis by Gumbel’s
Figure 2. Representative S-N curves for MH 65.17 sintered at 1,280°C and ABC 100.30 sintered at 1,120°C; sintered density 6.78 g/cm3
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TABLE I. EXPERIMENTAL RESULTS Sintering Temperature 1,120°C Powder
Density (g/cm3)
Fatigue Strength (N/mm2)
Hardness (HBW)
λ (µm)
δ (µm)
F (λ)
H/λ1/6
4πA* P2
ρ2 4πA P2
ABC 100.30
6.34 6.78 7.121 7.01 7.381,2 7.203
42.3 61.4 67.1 70.6 85.0 103.3
35 43 50 50 60 63
229.6 141.3 105.0 121.1 108.8 94.4
33.12 22.12 15.93 20.80 29.43 14.54
0.0407 0.0685 0.0840 0.0814 0.0671 0.1126
14.1 18.8 23.0 22.5 27.5 29.5
0.0733 0.1055 0.1430 0.1200 0.1122 0.1548
10.9 14.9 19.2 17.0 18.2 20.4
NC 100.24
5.854 6.064 6.364 6.674 6.941,4
45.0 59.5 74.6 82.9 89.0
42 44 53 61 71
184.5 161.5 123.4 110.1 92.2
34.23 27.43 22.63 17.94 26.27
0.0386 0.0471 0.0602 0.0879 0.1135
17.6 18.9 23.8 27.9 33.4
0.0833 0.0952 0.1138 0.1355 0.1509
9.9 11.3 13.6 16.4 18.7
MH 65.17
5.50 6.06 6.35 6.73
35.4 43.6 61.3 72.3
28 39 47 54
218.7 173.7 139.4 124.1
37.30 35.74 26.05 33.94
0.0377 0.0590 0.1024 0.1287
11.4 16.5 20.6 24.2
0.0838 0.1012 0.1485 0.1599
8.8 11.7 15.5 18.1
SINTERING TEMPERATURE 1,280°C Powder
Density (g/cm3)
Fatigue Strength (N/mm2)
Hardness (HBW)
λ (µm)
δ (µm)
F (λ)
H/λ1/6
4πA* P2
ρ2 4πA P2
ABC 100.30
6.41 6.83 7.151 7.12 7.391,2 7.203
54.1 71.9 94.6 96.6 117.3 105.8
39 47 59 58 67 58
202.4 139.6 89.1 107.7 75.1 89.8
36.03 24.72 14.40 26.81 19.54 13.84
0.0560 0.0781 0.1150 0.1006 0.1437 0.1161
16.1 20.6 27.9 26.6 32.6 27.4
0.0876 0.1248 0.1774 0.1472 0.1952 0.1669
12.2 16.5 21.5 19.5 24.1 21.2
NC 100.24
5.85 6.07 6.40 6.69 6.971
46.5 58.9 68.6 83.3 91.9
36 40 45 54 66
184.1 145.1 118.5 99.8 90.3
32.83 25.31 18.16 21.77 19.76
0.0419 0.0557 0.0815 0.0953 0.1118
15.1 17.4 20.3 25.1 31.2
0.0896 0.1004 0.1215 0.1466 0.1766
10.2 11.7 14.3 17.1 20.4
MH 65.17
5.56 6.02 6.41 6.78
49.8 60.2 79.9 100.7
32 45 53 62
198.5 162.3 124.2 111.9
31.63 30.20 26.60 31.46
0.0469 0.0709 0.1237 0.1441
13.3 19.3 23.7 28.2
0.0932 0.1235 0.1683 0.2039
9.4 12.7 16.9 20.8
1 Material delaminated
3 Double pressed/double sintered
2 Warm compacted
4 Sintered in 95 v/o N +5 v/o H +100 L/h CH 2 2 4
statistics of extremes. The pore size was taken as the largest dimension between parallel tangents to the pore (Figure 1), called the maximum Feret Fmax or caliper size of the pore. In addition, for all pores, the form or shape factor F = 4πA/P2 was calculated, where, according to Figure 1, A is the projected pore area and P is the pore perimeter. F approaches 0 if the pore degenerates to a line, and becomes 1 for circular pores. As shown earlier by Beiss and Dalgic6 there is a strong statistiVolume 45, Issue 2, 2009 International Journal of Powder Metallurgy
*Average of 2% most irregular pores
cal correlation between the shape factor F and caliper pore size Fmax which can be described by a Weibull distribution with a negative exponent. Statistically large pores in iron and steel are extremely irregular, expressed by a shape factor F < 0.2. The largest fifty pores were ranked in order of increasing size to calculate the probability of occurrence P = (i – 0.3) /(n + 0.4), where i is the rank and n (= 50) is the sample size. From a ranking of pore irregularity, the 2% of the most irregu-
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lar pores were selected to calculate their average shape factor. RESULTS Experimental results for the two sintering temperatures are included in Table I. Representative S-N curves of ABC 100.30 and MH 65.17 at a sintered density of 6.78 g/cm3 are shown in Figure 2. The irregular sponge iron sintered at 1,280°C exhibits superior strength in comparison with the water-atomized iron sintered at 1,120°C. Table I lists the endurance-limit amplitudes at the 50% probability survival level for all the experimental conditions, and the other test results. Gumbel Distribution The probability density function of Gumbel’s distribution for the quantity x is given by:
(
)
1 x–λ x–λ f(x) = — exp – exp –—— ·exp – —— δ δ δ
(1)
In the following, x is the size Fmax of the largest pores selected, λ is a characteristic pore size, and δ is a measure of the scatter or the width of the distribution. Figure 3 illustrates the shape of the distribution curve for differing values of δ and constant λ. The latter defines the location of the maximum, and increasing δ spreads the distribution. The right-hand end of the curve is stretched to high values of x. Figure 4 shows the effect of varying λ at constant δ. Increasing λ shifts the distribution curve from left to right. The cumula-
Figure 3. Probability density of Gumbel’s distribution at constant λ and differing values of δ
42
tive frequency function F(x) is obtained by integration of equation (1):
(
x–λ f(x) = exp – exp – —— δ
)
(2)
Figure 5. gives the five curves from Figures 3 and 4 as cumulative frequencies. Replacing F(x) by the estimated probability P = (i - 0.3)/(n + 0.4) and taking twice the logarithm results in: 1 l –ln(– ln P) = — x – — δ δ
(3)
Equation (3) gives a linear dependence between the pore size x (= Fmax) and the term –ln (–ln P). The coefficients λ and δ, parameters of the distribution, can be determined by linear regression analysis, Figure 6. Corresponding evaluations are presented in Figures 7 and 8 for the most irregular pores in ABC 100.30 sintered at 1,120°C and the most circular pores in MH 65.17 sintered at 1,280°C. For the linear regression, only the first 47 pores were included, because the largest 3 pores frequently deviate from the least square line and alter the slope 1/δ. In principle, the dimensions of the largest pores can be described by the Gumbel distribution. In the low-density versions of ABC 100.30, the large pores follow a bimodal distribution. There is a tendency for the slope 1/δ to become steeper at high densities. The characteristic pore size λ has a pronounced density dependence and
Figure 4. Probability density of Gumbel’s distribution at constant δ and differing values of λ Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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Figure 5. Cumulative frequency of curves in Figures 3 and 4
can be extrapolated to 0 at the pore-free density. At lower densities, λ is significantly smaller for the sponge irons than for the water-atomized irons. At densities ~7.0 g/cm3 the λ values converge, however, and no longer permit a clear distinction. Delaminated microstructures, for example, in the warm compacted ABC 100.30 iron sintered at 1,120°C, deviate from the general pattern with a steeper slope and smaller pores with increasing density. The unetched light optical micrograph (LM), Figure 9, of the microstructure explains why this is so. Delaminations can combine neighboring pores to create exceptionally large and irregular voids which account for the slope irregularities. Because of the difficulties associated with the three largest pores and the lack of a
Figure 7. Pore distribution of ABC 100.30 sintered at 1,120°C
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
Figure 6. Linearization of the cumulative frequencies in Figure 5
systematic reproducibility of the slope at higher densities, it was decided not to use the largest pore, as Murakami did,1 to estimate the fatigue strength. Instead, the characteristic pore size λ was used. This is justified as the size of the specimens was kept constant in this study and λ is virtually unaffected by slope changes caused by a single large pore. Form or Shape Factor The geometrical irregularity of plane objects in quantitative metallography is usually expressed by form or shape factors, the most common being F = 4πA/P2. It is generally assumed that highly irregular pores cause an internal stress concentration under mechanical loading which is detrimental to
Figure 8. Pore distribution in MH 65.17 sintered at 1,280°C
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fatigue properties. In most studies which attempted to take the pore shape into account to explain fatigue behavior, averages were calculated over all pores. As has been shown, however, about 50% of all pores are <5 µm and are essentially round which gives averages that are influenced only moderately by the base powder, density, or sintering temperature.11 It is not realistic to estimate the internal stress concentration from average values over all pores. For pore sizes >5 µm, the largest dimension Fmax essentially exhibits a logarithmic normal distribution.6 As a consequence, there is a long tail of a few large pores in the distribution on the right-hand side. Also, these pores are the most irregular ones whose shape factor can be estimated from the relation:
( )
4πA Fmax —— = 1– exp – ——— P2 a
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rounder than in the water-atomized iron sintered at 1,120°C. If the shape factor for each experimental condition is determined from the 50 largest pores for the characteristic pore size λ, there is virtually no deviation from the corresponding curves. For this reason the shape factor F (λ) from Table I is included in Figure 10. On average, the shape factor F (λ) is about 25% higher after high-temperature sintering than for conventional sintering. Surprisingly, this is consistent with a decrease in the characteristic pore size λ after high-temperature sintering of the water-atomized powder by about 15%, and for the sponge iron powder by about 10%. This takes place in spite of the disappearance of small pores and growth of larger pores at high temperatures. Figure 11
b
(4)
if a and b are known. The agreement with measured values is excellent if the number of pores available for evaluation is sufficiently large.6 Here, however, we are dealing with only 50 individual pores per experimental condition among which the scatter is high and not too systematic. Therefore, averages were calculated for each ten consecutive pores after size classification. These averages are shown in Figure 10 for ABC 100.30 sintered at 1,120°C and MH 65.17 sintered at 1,280°C as the two extremes in this study. The data scatter is high in the case of the sponge iron. Comparing equal pore sizes, the pores in the high-temperature-sintered sponge iron are much
Figure 10. Shape factor for the most irregular and the most circular pores, average of ten consecutive pores after size classification
Figure 9. Delaminations at a high density; sintering temperature 1,120°C; unetched. LM
Figure 11. Differences in pore morphology after conventional and high-temperature sintering
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attempts to explain this phenomenon. The growth of the large pores, at the expense of the small pores, is accompanied by pore rounding to an extent that compensates the asperities and gives smaller Fmax values at a larger pore area. This aspect is particularly pronounced with delaminated microstructures. During high-temperature sintering the diffusional mass flux is large enough to heal many delaminations and cause them to disappear by closure. Thus, they lose their adverse effect on mechanical properties. For comparisons the average shape factor of the 2% most irregular pores is listed in Table I. Since the number of pores contributing to this average is between 1,800 and 8,800, the mean shape factor is larger than F(λ). MECHANICAL PROPERTIES Sintered density is a primary variable governing mechanical properties. Figure 12 shows the fully reversed bending fatigue data from Table I as a function of sintered density. Material sintered at 1,120°C is coded with open symbols and the hightemperature-sintered irons coded with filled symbols. Delaminated microstructures are marked with an asterisk. Only about 80% of the data fit in a range of ±20% around a calculated average. The lower-strength results reflect predominantly the water-atomized powder sintered at 1,120°C. The higher end comprises mainly the pure broken sponge iron sintered at 1,280°C, as anticipated from Figure 2. The high-temperature-sintered delaminated specimens are not inferior to the defect-free microstructures. The variation between the highest and the lowest fatigue strengths (endurance limits) at equal sintered density is attributed to microstructural differences, neglecting experimental error. The density dependence alone is not sufficient to understand the fatigue strength of the sintered irons. From Table I it is clear that the water-atomized iron has a relatively low fatigue strength and a lower hardness than the sponge iron at the same sintered density. If the fatigue strength is plotted versus apparent hardness, a much better correlation is found, Figure 13. About 80% of the results are found within ±10% around a least square fitted average curve. The largest individual deviation is observed for a low-temperature-sintered delaminated test series. This confirms that delaminations cannot be detected by apparent hardness measurements because the sintered density– Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
Figure 12. Sintered density dependence of fully reversed bending-endurance limit (from Table I)
Figure 13. Effect of apparent hardness on fully reversed bending-endurance limit (from Table I)
hardness relationship is not affected by the small separations between adjacent powder particles perpendicular to the pressing direction. In contrast, the mechanical properties in tensile tests, fracture toughness tests, or fatigue tests respond sensitively to this type of defect. A similar conclusion was drawn by Ganesan et al.12 High-temperature-sintered materials exhibit statistically higher fatigue strength at equal hardness, and the negative effects of delaminations have essentially vanished. The apparent hardness, which is influenced by sintered density, is better correlated with fatigue performance than sintered density.
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Figure 14 shows the relationship between the endurance limit and apparent hardness divided by the sixth root of the characteristic pore size, H/λ1/6. The influence of hardness is dominant, because all the data falling out of the ±10% range in Figure 13 do so in Figure 14. Modifying hardness by the microstructural aspect, as suggested by Murakami,1 from fracture mechanics reasoning yields a linear function and slightly less scatter than hardness alone. However, the predictive accuracy has barley improved. In earlier publications6,11 a good correlation was established between the fatigue strength of
Figure 14. Relationship between fully reversed bending-endurance limit and H/λ1/6
unalloyed sintered iron and the square root of the shape factor of the 2% most irregular pores: σA ~ ρ2 4πA/P2
where ρ is the sintered density of the material. A corresponding plot is presented in Figure 15 based on the evaluations made during this work, the main difference from Beiss and Dalgic 6,11 being the size of the area analyzed and, thus, the number of pores. In the current work the number of pores was higher by about one order of magnitude. Therefore, the image analysis results here are more representative. Comparing Figure 14 and Figure 15, there is no significant difference in scatter or any advantage in favor of either one. Both relationships seem equally well suited for describing the interdependence between large pores and fatigue strength. DISCUSSION The fact that apparent hardness divided by a pore-size-related quantity, and sintered density multiplied by a shape factor, yield similar insight into fatigue behavior is not a contradiction, provided the microstructural characteristics are restricted to large pores which are capable of acting as crack-initiating stress raisers. Hardness and sintered density are closely related, and according to equation (4) the shape factor F is not independent of the pore size Fmax, as was shown in Figure 10. The scatter in Figure 10 may not be too convincing, but it must be appreciated that the data points represent averages of only ten pores. If the number of pores increases, the averages fall onto the descriptive curve with little variation.6 Thus, the main difference between the two approaches lies in the selection of the large pores. Therefore, a plot of fully reversed bendingendurance limit versus the product of hardness and shape factor of the 2% most irregular pores must yield a reasonable correlation, which is confirmed by Figure 16. Murakami’s1 original equation to estimate the fatigue strength is: σA=143(HV+120)A-1/12
Figure 15. Relationship between fully reversed bending-endurance limit and ρ2 4πA/P2, where is the average of the 2% most irregular pores
46
(5)
(6)
where HV is the hardness in the pore-free state, not the apparent hardness of a porous structure, and A is the projected area of the largest defect in a highly stressed volume. The characteristic large pore size λ is considered more representative of a Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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where B and C are constants deter mined by regression analysis. From Figure 17, B = 1.55 × 10-8 and C = 5.8. The result of this analysis is given in Figure 17. More than 80% of the experimental data fit a scatter band of ±15% about the non-linear average. The original for m of Murakami’s equation describes a trend, but is not appropriate for estimating the fully reversed bending-endurance limit of unalloyed porous iron.
Figure 16. Relationship between fully reversed bending-endurance limit and H (4πA/P2)
Figure 17. Relationship between fully reversed bending-endurance limit and hardness divided by the sixth root of the characteristic large pore size λ
microstructure than the random largest individual defect. Therefore, A-1/12 is replaced by λ-1/6. The hardness of pore-free pure iron can be estimated as 104 HV. In a regression analysis, aside from scatter, a slight curvature of σA is observed when plotted versus (104 HV + 120) / λ1/6. This deviation from linearity can be described by a power law. Neither the exponent of the power law nor the coefficient of correlation is changed when the hardness correction of 120 is removed from equation (6). The equation can be rewritten in the form: σA = B(104HV/λ1/6)C Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
(7)
SUMMARY AND CONCLUSION Gross structural defects such as exceptionally large pores must affect the fatigue strength of porous sintered materials from a fracture mechanics perspective. An extensive fatigue-testing program was undertaken with nominally pure iron powders to find out what microstructural features govern performance if alloying effects are excluded. One approach is to identify large defects by Gumbel’s statistics of extremes; the other possibility is the selection of a certain percentage of pores from the large-size tail of the pore-size distribution. Irrespective of choice, all large pores in die-compacted-and-sintered iron and steel are geometrically irregular. This irregularity is reduced during high-temperature sintering, and even delaminations can be healed to a significant extent. Sintered density alone is not sufficient to describe the fatigue strength of pure iron. Taking into account the relation between hardness and fatigue strength and microstructural features, scatter in experimental data can be reduced by a factor of 2 and reliable predictive equations can be derived. The strong correlation between hardness and fatigue strength implies that it is more important to improve hardness than to remove extremely large pores. REFERENCES 1. Y. Murakami, Metal Fatigue: Effects of Small Defects and Nonmetallic Inclusions, 2002, Elsevier, Amsterdam, The Netherlands. 2. E.J. Gumbel, Statistics of Extremes, 1957, Columbia University Press, New York, NY. 3. L. Ledoux and C. Prioul, “The Influence of Pore Morphology on the Monotonic and Cyclic Properties of Sintered Iron”, Moder n Developments in Powder Metallurgy, compiled by P.U. Gummeson and D.A. Gustafson, Metal Powder Industries Federation, Princeton, NJ, 1988, vol. 21, pp. 41–53. 4. L. Ledoux, “Effets de la porosité et d’un traitment de sur-
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5.
6.
7.
8.
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face—le traitment à la vapeur d’eau—sur les propriétés mecanique statique et cycliques du fer fritté”, 1989, thesis, École Centrale, Paris. T.M. Cimino, H.G. Rutz, A.H. Graham and T.F. Murphy, “The Effect of Microstructure on the Fatigue Properties of Ferrous P/M Materials”, Advances in Powder Metallurgy & Particulate Materials—1997, compiled by R.A. McKotch and R. Webb, Metal Powder Industries Federation, Princeton, NJ, vol. 2, part 13, pp. 137–149. P. Beiss and M. Dalgic, “Structure Property Relationships in Porous Sintered Steels”, Materials Chemistry and Physics, 2001, vol. 67, pp. 37–42. E.R. Rice and J. Tengzelius, “Die Filling Characteristics of Metal Powders”, Powder Metall., 1986, vol. 29, no. 3, pp. 183–194. G.F. Bocchini, “Influence of Small Die Width on Filling and Compacting Densities”, Powder Metall., 1987, vol. 30,
no. 4, pp. 261-266. 9. Y. Ozaki, S. Uenosono, N. Tagami, K. Kuwagi, R. Noda and M. Horio, “Experimental and Numerical Investigations of the Filling of Iron Powders”, Advances in Powder Metallurgy & Particulate Materials—2006, compiled by W.R. Gasbarre and J.W. von Arx, Metal Powder Industries Federation, Princeton, NJ, vol. 1, part 3, pp. 35–46. 10. U. Engström, P. Hofecker, D. Milligan and D. O’Keefe, “Utilizing Bonded Mix Technology to Improve the Dimensional Consistency of Complex, Tightly Toleranced PM Components”, ibid., pp. 54–61. 11. M. Dalgic and P. Beiss, “Einfluss des Gefüges auf die Schwingfestigkeit von Sintereisen und -stahl”, Prakt. Metallographie, 1999, Sonderband, vol. 30, pp. 97–105. 12. P. Ganesan, S¸. Doms¸a and P. Beiss, “Fracture Toughness of PM Alloy Steels”, Powder Metall., 2005, vol. 48, no. 4, pp. 323–328. ijpm
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PM METALLOGRAPHY
EVALUATION OF PM FRACTURE SURFACES USING QUANTITATIVE FRACTOGRAPHY Thomas F. Murphy, FAPMI*
INTRODUCTION Materials fail when exposed to stresses greater than their characteristic strengths. Whether the stresses are tensile, compressive, impact, or fatigue, exceeding the material strength limits will produce failure, frequently accompanied by the creation of a fracture surface. When PM materials fail, the fracture surface is different in appearance when compared with failures from more traditional materials. These differences include the smooth surfaces of the inherent porosity and localized variations in fracture caused by differences in the microstructure. PM microstructures may be heterogeneous and complex due to the distribution of the alloying additives, local hardenability, and part processing. With most PM parts, the majority of the porosity exists as a complicated interconnected network and the microstructure ranges from uniform to an aggregate of transformation products and phases.1–8 The alloying method used in the manufacture of the iron-base powders, coupled with the part density and sintering conditions, determine the amount and location of the transformation products and how the part responds on loading. All of these variables contribute to the properties of a finished PM part and the appearance of a fracture surface if a part fails. Therefore, analyzing and quantifying the features and geometric characteristics of cracks and fracture surfaces can provide invaluable information about the performance of PM materials under stress. In the manufacture of PM parts, the initial processing step is compaction. During compaction, the base powders and alloying additives are forced into intimate contact within a fixed, enclosed die. It is at this stage that the green density is established and the initial character of the pore network formed. Upon sintering, alloying of the additive materials occurs through diffusion and the contact regions between particles form metallurgical bonds known as sinter necks which strengthen the compact into a coherent part. The characteristics of the base powders, internal lubrication, and compaction pressure define the density. Figure 1 illustrates the appearance of the pore network inside a compact.1–3 It is an image of an epoxy skeleton formed by vacuum impregnating a porous sintered iron bar with epoxy, followed by
Fracture surfaces produced by breaking powder metallurgy (PM) materials appear substantially different when compared with those generated on parts fabricated by other metalworking techniques. Although the characteristics of the fractured regions are the same (ductile dimples, cleavage, etc.), the presence of interior pore surfaces and heterogeneous microstructures add complications to a fractographic analysis. Techniques are described to evaluate cracks and fracture surfaces in PM materials. Scanning electron microscope (SEM) analysis of three-dimensional (3D) surfaces is used to quantify the amount of fracture on the exposed surfaces and light microscopy (LM) on metallographically prepared cross sections is used to evaluate and quantify the location of cracks in relation to the microstructure and the roughness of the fracture profiles.
*Scientist, Research and Development, Hoeganaes Corporation, 1001 Taylors Lane, Cinnaminson, New Jersey 08077-2017; E-mail:
[email protected]
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extraction of the iron from the epoxy/iron composite using dilute acids. After removing the metallic portion of the composite, the remaining epoxy is a replica of the porosity. It shows the former locations of the iron particles as the large vacated areas within the structure. The gaps in the resin are the prior sites of metallic contact where sintering necks were created. It is clear from the shape in the central portion of the figure that an iron particle occupied this space and the holes in the epoxy are the necks formed with neighboring iron particles. These necks are often the weakest areas in the cross section and possibly the eventual location of fracture. To counter this and strengthen the compact, the necks can be strengthened in multiple ways. Two commonly used techniques include increasing the part density, which results in an enlargement of the neck
Figure 1. Epoxy skeleton showing morphology of the pore network. SEM/Secondary Electron Image (SEI)
(a)
areas, and selective diffusion, such as liquidphase sintering, which can strengthen ferrite or increase the local hardenability. Separately or together, these methods will result in the development of larger neck regions and/or stronger transfor mation products in the necks. Consequently, with the increased local strength levels, particularly in the sinter necks, fracture can be transferred to weaker locations within the structure, which may be in areas other than the neck regions. The alloying method used to manufacture the ferrous base powder also has a significant effect on the response of a PM material under load due to the local chemical composition and the resulting local hardenability. Distribution of the alloying elements is first controlled by the method used to introduce them to the iron or steel powder. Figure 2 shows a comparison of two microstructures from parts having the same overall chemical composition (Fe-4 w/o Ni-1.5 w/o Cu-0.5 w/o Mo-0.6 w/o graphite), but produced from base powders made using different manufacturing methods. In image 2(a) all the alloying additives (except graphite) were added to the melt (prealloyed) producing a uniform chemical composition between and throughout each particle. In image 2(b), the nickel and copper were partially alloyed to a 0.5 w/o Mo-base prealloyed steel during an annealing operation (diffusion-alloyed) resulting in a heterogeneous chemical composition within each particle. Due to the differences in local chemical composition, the microstructure produced with the diffusion-alloyed powders is similar to that found in compacts made from admixed powders.
(b)
Figure 2. Comparison of microstructures two PM steels with the same overall chemical composition but manufactured using different methods: (a) prealloyed, and (b) diffusion-alloyed. Etchant: 2 v/o nital + 4 w/o picral
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Further, the variability in local chemical composition, as seen in the diffusion-alloyed material, normally is not homogenized during standard sintering operations. The resulting microstructure is a combination of transfor mation products throughout the compact. Because of these differences in local chemical composition and microstructure, the two materials shown in Figure 2 will behave differently when loaded and are expected to produce different fracture surface morphologies/appearance. Upon failure and creation of a fracture surface in a PM part, the characteristics of the individual fractured areas are the same as those found in parts made with more conventional materials.9 However, while the individual fractured regions are the same, two major differences are apparent. The first is the smooth appearance of the interconnected pore network inherent to PM materials, and the second is the combination of fracture types with their locations, which may be caused by a heterogeneous microstructure. It should be emphasized that the extent of fracture on the surface is the effective load-bearing cross section of the part. The characteristics on the surfaces can be analyzed using quantitative fractography. Three quantitative techniques are used to describe the following: • amount of fracture on the projected 3D fracture surface • roughness of the fracture surface profile • location of enclosed cracks, especially in fatigue specimens, and their relationship to the microstructure In examining the fracture surfaces and cracks, it must be noted that the scale of the details ranges in size from microscopic to macroscopic.10 The analysis of these features uses the microscopy techniques known as fractography and, in cases where stereology or image analysis techniques are combined with fractography, quantitative fractography. Quantitative fractography uses rigorously defined stereological methods to produce quantitative information on the geometric characteristics and microstructural features present on the exposed and enclosed surfaces. 11–22 Parameters calculated include number, length, area, percent, etc. These stereological procedures can be used in the analysis of a wide variety of failed parts and test pieces. From in-service failures to the evaluation of research materials and parts, quantitative fractography can Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
add significantly to the understanding of material behavior under stress. METALLOGRAPHIC SAMPLE PREPARATION Fracture surfaces are delicate and require careful, well-planned handling to preserve the fine microstructural details.23 In many cases, these details are the only evidence surviving the failure that can be used to explain the material behavior under stress. As a result, following correct sample preparation procedures is critical in maintaining any and all the evidence of failure. Three analytical techniques are used in this paper to evaluate the geometric shapes and locations of cracks and separated surfaces. While normal preparation procedures are sufficient for some sample types, others require the use of special preparation routines. The exposed surfaces used for SEM examination often require little additional sample preparation other than careful cleaning. In comparison, cross sections may require the use of special sectioning, grinding, or polishing techniques in addition to the deposition of protective coatings. The areas where cross-sectional analysis is preformed are selected in two ways. First, as carefully chosen areas containing enclosed cracks, at spall failures or fatigued regions. Second, as vertical sections through the exposed fracture surface where the characteristics of the fracture profile can be evaluated. Often, samples containing enclosed cracks may be prepared using standard metallographic procedures, although depending on the extent of cracking, epoxy impregnation may be useful in maintaining the spatial relationships of the metallic fragments within the cracked regions. Vibratory polishing is also useful in revealing an accurate representation of the internal, enclosed cracks. Preparation of the vertical sections requires meticulous handling to avoid the introduction of artifacts and the destruction of details. This usually requires the application of a protective coating over the exposed surface. Fortunately, after an effective protection layer has been applied, normal metallographic preparation procedures can be used. Once the surface has been protected and prepared, the fracture surface edge is clearly delineated and accurate measurements can be made. Vertical sections are usually removed using abrasive cut-off saws; however, extreme care must be exercised not to introduce damage to the surface during cutting. Wafering saws equipped with thin abrasive or diamond blades are often recom-
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metallic specimen and the dark mounting material, thoroughly covering the surface. Consequently, the fine details are maintained, while providing a hard edge to help maintain a planar surface. QUANTITATIVE FRACTOGRAPHY A combination of SEM analysis on the area of fracture and LM on prepared cross sections provides invaluable information on the behavior of materials.
Figure 3. Stainless steel impact bar with fracture surface protected with an electroless nickel coating. LM/unetched
mended as a preferred alternative. After sectioning, a compression mounting system can be used to mount unprotected cross sections, but with the possibility of major drawbacks. Due to the pressure required for the formation of the mount, the compression mounting of powders could increase the risk of deformation to the fracture surface through transfer of pressure to the small surface details. In addition, the hard fillers included in the mounting powders to maintain surface flatness can agglomerate along the fracture edge, preventing the mounting material from creating precise contact with the surface and leaving a gap between the sample edge and the mounting material. The best and most reliable method for protecting the sample edges before compression mounting appears to be one of the electroless nickel-plating systems, in which a hard nickelphosphorus layer is deposited on the surface.24 Obtaining correct results with these systems relies on close adherence to procedural detail, especially in maintaining the correct plating temperature. The process involves a thorough cleaning of the sample, plating the fracture surface and the surrounding area, and allowing the sample to dry. Once dry, a vertical section can be removed using a wafering saw. In sectioning, the direction of the cut must be into the fracture surface rather than away from it to prevent peeling of the plating from the surface. Figure 3 shows the result of electroless nickel plating on a high-temperaturesintered stainless steel impact fracture. This sample was used to demonstrate the effectiveness of the coating on a highly jagged surface. The protective plating in seen between the bright, porous
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Area Percent Fracture Procedure As with most analyses, correct sampling of the separated fracture surface is required to provide confidence and accuracy in the evaluation. The location of the fields used for measurement must be carefully chosen to provide an overall estimate of the amount of fracture on the surface, or areas must be selected in advance to give information on specific regions. In the first case, a systematic sampling scheme should be designed before the start of testing. In many cases, if the fracture surface is handled with care, careful cleaning may be the only preparation step required prior to SEM examination. Determining the area fraction (AA) of the features present in the SEM image can be performed using manual or automated methods.25–27 With the PM surfaces, a simple point count using an overlay of an x/y grid can give a reliable assessment of the amount of fracture and pore structure present. However, it must be noted that the SEM image does not represent the entire 3D surface, but is a projection of the surface. Regions oriented in directions deviating from the horizontal plane of view are not imaged in the same proportion as the horizontal areas. With this in mind, the results are reasonable estimates of the area fraction being measured on the projection, but not an estimate of the entire fracture surface because the entire surface is not uniformly sampled. The use of projections for measurement requires a change in the symbols used to designate the relationship of the measurement. A A (area of features of interest divided by test area) is changed to A'A, with the prime character indicating that the measurement was made on a projected image. The procedure for performing the count, a systematic manual point count, is well documented and easily accomplished.25–27 Live SEM images or photomicrographs of individual fields can be used as the subject material. An equally spaced x/y Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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Figure 4. Fractograph with x/y grid overlay. The 63 points at intersections of the x and y lines are the sampling locations. In this example, 18 of the 63 total points (delineated by arrows), coincide with fractured areas (~29% of sampled area). SEM/SEI
grid is overlaid on the image and counts are made when the x/y crossing points correspond to the features of interest. The counts are accumulated for multiple fields and this sum is then divided by the total number of grid points applied to the images. The result is an estimate of the area fraction of the features of interest over the total surface area. Multiple fields must be measured and the locations of the fields chosen on a random or, more preferably, a systematic basis. If specific areas are in question, entire areas or surfaces can be sampled and measured. The calculated value is dimensionless and the test method is accurate for any selected magnification. In developing the test procedures, careful consideration should be give to the magnification used. The magnification must ensure that all features have an equal opportunity for sampling. Introducing a bias due to feature size or shape should be avoided as much as possible. Figure 4 shows an example of a single field from a fractured PM specimen with an overlaid x/y grid. The literature discussing analysis of projected images should be consulted before designing a test procedure.14,17,18,21,25–27 As an additional benefit, the SEM imaging of the fracture surface also provides an opportunity to make a quantitative assessment regarding the types of fracture present (ductile, brittle, etc.). Materials and Processing The combination of part density, microstructure, and the method of fracture determine the Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
characteristics present on the fracture surface. In this example, a mixture of iron powder (wateratomized), copper, and graphite (FC-0205) was used as the test material to illustrate how the processing variables affect the appearance of the fracture surface and also to demonstrate the use of the metallographic methods. Transverse rupture (TR) bars were pressed to three green densities, 6.5, 6.8, and 7.1 g/cm 3 and sintered at 1,066°C (1,950°F), 1,093°C (2,000°F), and 1,120°C (2,050°F) in a belt furnace at a constant belt speed of 7 cm/min under a 90 v/o N2/10 v/o H2 atmosphere. This resulted in a sintering time of approximately 15 min at temperature, with the time at temperature defined as a temperature ±5°C of the intended sintering temperature. The density of the bars from the three green-density groups was measured after sintering. The 6.5 g/cm3 density green bars ranged in sintered density from 6.40 to 6.45 g/cm3, the 6.8 g/cm3 density green bars ranged from 6.65 to 6.73 g/cm3, and the 7.1 g/cm3 density green bars ranged from 6.96 to 7.04 g/cm3. Physical and mechanical properties were measured on the bars. The broken bars were examined using both LM and SEM to determine differences in the details on the fracture surface in addition to typical cross-sectional appearances.28 A'A Measurements Live SEM images of the fracture surface were analyzed for A'A using the x/y grid overlay.3,29 It should be noted that the type of fracture was irrelevant in this evaluation and any fracture present was counted where appropriate. Figure 5
Figure 5. Relationship of projected area percent fracture and sintered density with sintering temperature
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TABLE I. COMPARISON OF POROSITY WITHIN THE COMPACT WITH PROJECTED AREA MEASUREMENTS Density (g/cm3) Volume % Porosity in Compact Projected Area % Porosity Projected Area % Fracture
6.5 17 76–83 17–24
6.8 13 70–76 24–30
7.1 10 60–68 32–40
shows the results of the systematic point counts. As expected, the lowest densities at each sintering temperature produced the least material fracture with the amount of fracture increasing at a relatively constant rate with an increase in density at each sintering temperature. It should be noted that, at the three densities used, the volume percent porosity for each density was approximately 17 v/o for 6.5 g/cm3, 13 v/o for the 6.8 g/cm3 samples, and 10 v/o at 7.1 g/cm3. In Table I, a comparison of the volume percent porosity in the compacts is made with the amount of fracture and pore surface on the fracture surfaces. This large discrepancy in comparing the volume percent porosity within the compact with the amount of pore surface on the fracture is caused by the process used to create the fracture surface. A randomly selected plane through a sintered part
reveals approximately the same fraction of porosity as is contained in the material volume. However, the fracture surface produced by overstressing a part is not a random surface, but one created by the relative strengths of localized areas under stress. The areas of fracture on the projected surface represent specific weak regions in the load-bearing cross section and not randomly selected sites in the microstructure. The cracks tend to follow the pores because the weakest sintered regions are often the smallest contact areas. Alloying during the sintering process can change the individual locations of fracture by strengthening localized areas. These are often the sinter necks, which can be the weakest regions in the sintered compact. With a strengthening of the sinter necks through increases in ferrite strength or local hardenability, the load could be transferred from the necks to transparticle areas. Samples sintered at the two highest temperatures (1,093°C and 1,121°C), were strengthened in this manner. At these temperatures, which are above the melting point of copper, the liquid copper ran along the pore surfaces and diffused into the sinter necks and the interconnected porosity. Figure 6 shows differences in the fracture characteristics
Figure 6. Comparison of fracture characteristics in bars at 7.1 g/cm3 density sintered at (a) 1,066°C and (b) 1,120°C. SEM/SEI
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at sintering temperatures both below and above the melting point of copper (1,083°C) with compacts pressed to the same green density (7.1 g/cm3). Image 6(a) was sintered at 1,066°C and shows fracture located only in neck regions. In contrast, image 6(b) was sintered at 1,121°C and shows a combination of transparticle fracture and fracture at neck regions. This change in fracture location alters the amount of fracture on the surface and the characteristics of the fracture. It also provides evidence of a change in mechanical and physical properties, and, consequently, material behavior. Surface-Roughness Measurements The main difficulty in quantifying the features on a fracture surface is very basic. In order to develop relationships of the various features on the surface with the actual surface area, the area of the fracture must be known. Once the area has been measured or calculated, other quantities of interest can be determined. It is the measurement of the surface area of the fracture that has presented problems for many years. As was demonstrated with the A' A measurement, many useful evaluations are performed using SEM, although this area viewed does not represent the entire surface but is a projection of the surface. To overcome this problem, techniques have been devised to analyze planar, vertical cross sections taken through the fracture surface, where measurements are made of the fracture profile and then used to determine the actual fracture surface area. In this analysis, the importance of the metallographic sample preparation cannot be overstated. Making profile measurements requires utilization of the best possible metallographic sample preparation techniques to assure accuracy in the representation of the details along the sample edge. The type of stress placed on the part determines the number and location of the sections required for this analysis. Appropriate literature should be consulted to ensure that the correct test procedures are followed for particular samples. These vertical sections are used to determine the roughness parameter of the fracture profile, R L, which is calculated by dividing the actual length of the fracture surface by its projected length, as dictated by equation (1):30–31 RL = λ / LP Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
(1)
where RL is the fracture profile roughness parameter, λ is the actual length of the fracture, and LP is the projected length. This value is dimensionless and ranges from 1 for a flat surface to a theoretical value approaching ∞ for an extremely irregular one. By itself, it can be used to compare surface profiles. However, it is more frequently used as an important component in calculating the fracture surface-roughness parameter, RS. For the image presented in Figure 3, the length of the interface between the nickel plating and the stainless steel sample is measured as the fracture surface profile length (λ) and the horizontal length of the image is the projected length (LP). Numerous methods using RL to estimate fracture surface area were developed in the 1980s with varying degrees of success. 12,16–19,32–37 Advances in the early attempts were made in the early 1990s, when an unbiased technique using a combination of the angular distribution of individual fracture-profile segments with R L was made.21,30,38,39 In this technique, the profile of the fracture is separated from the image and divided into small segments, as seen in Figure 7(b). The angle of each segment is measured and the frequency distribution of the angles determined using 18 10° classes. The value of each angle is measured with the vertical direction at 0°, as shown in Figure 7(a). Each angular class is numbered, starting at class 1 (the lowest angular value) and proceeding to class 18. The fraction of each class within the total distribution is then calculated. Equation (2) is the relation used to calculate
Figure 7. (a) location of the 18 10° angular classes. (b) portion of a fracture profile separated into the line segments. The direction of each segment is compared with the angular directions in (a) to locate its class
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TABLE II. COEFFICIENTS FOR CALCULATION OF THE PROFILE STRUCTURE FACTOR (Ψ)* Class I
Coefficient (a)
Class (i)
1 2 3 4 5 6 7 8 9
1.565 1.5232 1.4508 1.3599 1.2655 1.1694 1.0906 1.0336 1.0037
18 17 16 15 14 13 12 11 10
*Determined by the angular distribution of the fracture segments the fracture surface-roughness parameter, RS. It is the roughness-parameter calculation for the fracture surface and is the analogous calculation to RL, which is the roughness parameter for the fracture profile. RL is calculated by dividing the measured length of the surface profile by the projection of the profile. RS is a calculation of the area of the fracture surface divided by the projection of the total surface. Values of RS also range from 1 for a flat surface to approaching ∞ for a highly irregular surface. The value of Ψ is determined using equation (3) and the coefficients (ai) found in Table II. RS = RL • Ψ Ψ=
(2)
1
Σ ah 18
i i
(3)
where Ψ is the profile structure factor, ai is the coefficient from Table II, and hi is the fraction of each class in the total distribution. In calculating the value of RS, it can be seen in Table II the coefficient values are symmetric around the horizontal direction (90°). Since the minimum value for RS is 1, it follows that the coefficient around the 90° angle (classes 9 and 10) is the lowest and the value of RL will also be nearly 1. Further, as the angles diverge from 90° (classes 9 and 10), the coefficient values become larger, the RL becomes greater because the surface is more irregular, and RS will increase. Determination of RL and RS The FC-0205 TR samples used in the projected area fraction measurements were also used to illustrate the measurements of R L and R S .
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Samples from the three sintering temperatures at a sintered density of 7.1 g/cm3 were selected for quantitative assessment. The fracture surfaces were protected with an electroless nickel plating and standard metallographic techniques used for preparation. In order to facilitate testing and provide sample-to-sample uniformity, a program was developed for an automated image analysis (AIA) system to perform the RL and RS analysis on the prepared surfaces. This analysis was accomplished in the following stages: • separating the fracture edge from the plated surface • measuring the total fracture length and horizontal projection • dividing the fracture profile into line segments • determining the angular measurements of the segments of the fracture surface The data were exported into a Microsoft Excel spreadsheet where the frequency distributions and R S calculations were made. Additionally, cracks were introduced into one sample prior to sintering to create the appearance of green cracks. It was then sintered, broken, and prepared for comparison. Using the automated program, the entire fracture surface profile was sampled and measured. The results are displayed in Table III. In observing the data, the R S values of the green-crack surface and the sample sintered at 1,066°C are similar. With the low sintering temperature, the copper has not melted, alloying has not taken place, and the amount of sintering is relatively low. Possibly, the fracture paths in both cases are similar. Fractures occur in the green condition at locations where particles are forced into contact during compaction. These are probably at the smallest contact areas and would be the same, or similar, in the lowest-temperature-sintered sample because of the lack of additional alloying from the copper and minimal sintering. Upon melting of the copper at the 1,093°C sinter, the RS value drops, indicating a flattening of the TABLE III. RESULTS OF RL AND RS TESTING ON SINTERED (7.1 g/cm3) TR BARS AND BAR CONTAINING GREEN CRACKS Sample ID
RL
RS
Green Crack 1,066°C (1,950°F) 1,093°C (2,000°F) 1,120°C (2,050°F)
2.792 2.729 2.300 2.727
3.584 3.583 2.950 3.451
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surface and less reliance of failure through the smallest neck regions. With a further increase in sintering temperature to 1,120°C, RS increases almost to the value of the green-crack and 1,066°C samples. This may be due to further alloying of the copper and a change in the fracture characteristics, including an increased amount of cleavage. It is interesting to note that both the RL and RS values in Table III are substantially higher than those typically observed in other materials. Examples using different materials are: (i) fatigued specimens from cast irons D4018 (1.7 to 2.4) and D5506 (1.3 to 2.2),21 (ii) low-temperature impact fractures from low-carbon iron and 11Cr2Ni steel (1.6 to 1.9),20 and (iii) tensile fractures from an AlLi–Al2O3 fiber composite (1.83 to 2.74 in one study38 and 2.17 in another39). This confirms that the fracture topography of the PM materials contains significantly more irregularity than the other cited materials and probably is affected by the surfaces of the inherent interconnected pores and the inclusion of porosity into the fracture surface. Again, the fracture surface is the load-bearing cross section through the stressed region of the PM part. It is determined by a selection of the weakest regions in the local structure, which are small sinter necks, locations with large pores, pore clusters, and microstructurally weak areas. Crack Proportions As a material with a complex and/or porous microstructure is stressed, particularly if the loading is cyclic (fatigue), crack growth may proceed in a preferential manner. It may favor growth through selective microstructural constituents while avoiding others. Also, the growth may proceed from one feature to another, such as from pore to pore.40–48 Sometimes the progress of the crack is along specific angles or at precise depths below the sample sur face. In other cases, microstructural constituents may act to blunt the progress of the crack. Often, the growth is random through a part, apparently without regard for the feature locations or transformation products present in the microstructure. If a crack progresses along a random path through the microstructure, the proportions of the constituents containing crack segments should be equivalent to the volume fraction of those constituents within the microstructure in that vicinity. On the other hand, if the crack proVolume 45, Issue 2, 2009 International Journal of Powder Metallurgy
gresses along a preferred path, for example, through particular microstructural features such as a transformation product, then a higher proportion of the crack will be located within those features. Figure 8 illustrates this concept using a micrograph of an unetched PM material containing pores. Three hypothetical crack paths are shown through the microstructure. The top line is a random path, the middle line shows preferential travel through the pores, and the bottom line avoids the pores. The proportions through specific constituents are calculated from equation (4): P(x) = AA(x)(crack) / AA(x)(total)
(4)
where P(X) is the proportion of a crack within a specific microstructural constituent (x), AA(X)(crack) is the fraction of constituent (x) containing crack segments, and AA(X)(total) is the fraction of the sample area occupied by constituent (x). It should be emphasized that sampling should be as large as is practical (multiple fields) to gain as much accuracy as possible and the magnification should be carefully chosen to ensure accurate representation of the microstructural features present. Specific portions of the crack can be quantified if desired. This may be useful if slow fatigue-crack growth is to be analyzed, rather than cracking in the fast-fracture region. The results of the calculation are dimensionless and, if percentages are desired, the fractional measurements can be multiplied by 100. The magnification should be carefully chosen to ensure accuracy in sampling the microstructural constituents and cracks. An illustration of the use of this test method is shown using a cross section from a PM part sub-
Figure 8. Example of crack paths through a porous microstructure. The hypothetical cracks exhibit random and non-random growth from pore to pore. Unetched/LM
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jected to cyclic loading. Figure 9 shows the cracked area within the etched microstructure. The first step in determining if the crack has taken a path located within preferred or randomly selected regions is to evaluate the microstructural constituents within that specific area. An image analysis system was used to determine the area percentages for each microstructural constituent. In this simple case, there are only three: martensite, nickel-rich regions, and pores. Separation of both the pores and nickel-rich regions from the remainder of the microstructure is required to determine the area fraction/percent of each. This
is shown in Figure 10(a) where a gray-level threshold was set to match the gray shade of each set of features. The blue features are the pores and the purple are the nickel-rich areas. Area measurements were made on these two areas and the amount of martensite was deter mined through subtraction. The results are shown in Table V and will be used for the proportion comparison with the location of the cracks in the cross section. The values shown in Table IV are used as the basis for determining the randomness or preference of the crack location. Figure 10(b) shows the cracks highlighted with the red lines. The total length of the line segments was measured, separated by their location within each constituent, and the total length within each constituent determined. The results are shown in Table V as a TABLE IV. COMPOSITION OF MICROSTRUCTURE* Microstructural Constituent
Area (% Measured)
Martensite Nickel-Rich Regions Porosity
86.9 (by subtraction) 5.5 7.6
* Determined using automated image analysis TABLE V. COMPOSITION OF AREAS CONTAINING CRACKS*
Figure 9. Etched image of the cracked region of a fatigued PM part. The microstructure is primarily martensitic with nickel-rich areas and pores. LM, 2 v/o nital + 4 w/o picral
Microstructural Constituent
Area (% Measured)
Martensite Nickel-Rich Regions Porosity
66.2 5.4 28.4
* Determined using automated image analysis
Figure 10. (a) image shows gray-level separation of porosity (black) and nickel-rich areas (gray ) in the martensitic matrix (white). Cracks are shown in image (b) as black lines. LM
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percentage of the total crack length. Using the measurements from Tables IV and V, the P (x) values were calculated for each constituent to determine if a preferred path was taken as the crack progressed. Equations (5), (6), and (7) show the P(x) calculations for each constituent. Pmartensite = 66.2 / 86.9 = 0.76 (
(5)
PNi-rich = 5.4 / 5.5 = 0.98 (Approximately Random)
(6)
Ppores = 28.4 / 7.6 = 3.74 (>Random)
(7)
Consequently, the crack in this example moved on a path preferring a pore-to-pore progression, with the nickel-rich areas chosen on a random basis, and the martensite less than random. SUMMARY Techniques have been presented to analyze fracture surfaces and cracked regions in PM materials. SEM was used to quantify the amount of fracture on exposed surfaces and LM was used to evaluate both the fracture-profile roughness and the relationship of cracks to the microstructural constituents within the part. In combination, they provide the possibility of solving real-world issues with part properties by evaluating characteristics revealed through failure. Individually, they each provide a different view into the behavior of materials under various loading conditions. In combination, they complement each other, making the picture of part behavior under stress more complete and they improve the information-gathering capability of the researcher. In the area-fracture analysis, SEM techniques were used to estimate the amount of fracture on separated surfaces. It was shown how the density and sintering temperature affect the fracture, in addition to increasing the amount of alloying as the sintering temperature was raised. The strengthening of the neck regions through increased alloying and improved particle-to-particle bonding in the neck regions also resulted in a transfer of the load to weaker transparticle regions, with the result being an increase in actual fracture on the surface. The use of LM added another beneficial means
Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
of looking at the fracture surfaces. Unlike SEM analysis where supplemental sample preparation may not be needed, the techniques used to prepare samples for analysis using LM are extremely important. Protection of the surfaces and correct representation of the microstructure must be accomplished if the analysis is to be meaningful. With fracture surfaces, nickel plating is a reliable and reproducible technique used to both protect the surface and provide a hard barrier to prevent edge rounding. Calculations of surface roughness through measurements of fracture profiles and angular distributions of profile segments can be beneficial in the comparison of fracture surfaces from different materials or those made using different processes. The comparison of surface irregularity can be an indicator of how a material is strengthened and the combination of cross-sectional examinations in the unetched and etched conditions can provide clues as to how a material fails. As the sintering temperature is raised and/or the density increased, the surface characteristics change along with the areas providing the strength of the compact, especially with the introduction of a liquid phase such as copper in FC0205 materials. The areas prone to fracture prior to the introduction of the liquid copper now may not be the weakest areas and the crack profile becomes flatter, even though more fracture area is seen. This finding indicates a larger load-bearing cross-sectional area with an increase in the overall material strength. With increases in sintering temperature, the surface changes are probably attributable to the distribution of the copper and strengthening of the ferrite. These are characterized by increases in the material strength, the amount of brittle fracture, and in the total amount of load-bearing surface area. The third technique involves a quantification of preferential crack growth through specific microstructural constituents or from feature to feature. This is particularly useful in fatigueloaded samples where cracking proceeds slowly until a critical crack length is reached and the fracture proceeds rapidly. It showed how the proportions of the constituents containing crack segments could be estimated and compared with the proportions of the constituents in the overall microstructure. For example, if a crack proceeds in a random path through the microstructure, the
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proportions of the constituents containing crack segments should be equivalent to the volume percentage of those constituents in the microstructure. Conversely, if the crack proceeds in a preferred manner, the proportion of the microstructure containing crack lengths will be greater in the preferred constituent than the volume fraction of that constituent. In this case, other constituents will contain a less than random amount of the crack. The percentage of the microstructure containing cracks will vary, with the preferred constituent containing more of the crack lengths, while other constituents will contain lesser amounts. ACKNOWLEDGEMENTS The author would like to thank colleagues, both former and current, for their assistance in preparing this manuscript. Paul Kremus, a former colleague, was responsible for making and testing the FC-0205 sintered samples used in the surface evaluations. His thoroughness and attention to detail is greatly appreciated. Bruce Lindsley, a current colleague, has provided guidance and encouragement throughout the preparation of this paper. In addition, I was fortunate to receive comments and advice from Arun Gokhale, Georgia Institute of Technology, on quantitative fractography-related procedures and results.
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pp. 961–971. 20. X.W. Li, J.F. Tian, Y. Kang and Z.G. Wang, “Quantitative Analysis of Fracture Surface by Roughness and Fractal Method”, Scripta Metallurgica et Materialia, vol. 33, no. 5, 1995, pp. 803–809. 21. A.M. Gokhale, “Quantitative Fractography”, Failure Analysis and Prevention, ASM Handbook, ASM International, Materials Park, OH, 2002, vol. 11, pp. 538–556. 22. K. Slamecka, P. Ponizil and J. Pokluda, “Quantitative Fractography in Bending-Torsion Fatigue”, Materials Science and Engineering A, 2007, vol. 462, pp. 359–362. 23. R.D. Zipp, “Preparation and Preservation of Fracture Specimens”, Fractography, ASM Handbook, ASM International, Materials Park, OH, 1987, vol. 12, pp. 72–77. 24. G.F. Vander Voort, Metallography Principles and Practice, 1984, McGraw-Hill, Inc., New York, NY, pp. 538–539. 25. R.T. DeHoff and F.N. Rhines, Quantitative Microscopy, 1968, McGraw-Hill, Inc., New York, NY. 26. E.E. Underwood, Quantitative Stereology, 1970, AddisonWesley Publishing Co., Reading, MA. 27. K.J. Kurzydlowski and B. Ralph, The Quantitative Description of the Microstructure of Materials, 1995, CRC Press, Boca Raton, FL. 28. B. Lindsley, G. Fillari and T. Murphy, “Ef fect of Composition and Cooling Rate on Physical Properties and Microstructure of Prealloyed P/M Steels”, Advances in Powder Metallurgy & Particulate Materials—2005, compiled by C. Ruas and T.A. Tomlin, Metal Powder Industries Federation, Princeton, NJ, 2005, part 10, pp. 353–366. 29. M. Slesar, E. Dudrova and E. Rudnayova, “Plain Porosity as a Microstructural Characteristic of Sintered Materials”, Powder Metallurgy International, vol. 24, no. 2, 1992, pp. 232–237. 30. J.R. Pickens and J. Gurland, “Metallographic Characterization of Fracture Sur face Profiles on Sectioning Planes”, Proc. Fourth International Congress of Stereology, NBS Special Publication 431, edited by E.E. Underwood, R. de Wit and G.A. Moore, U.S. Department of Commerce, National Bureau of Standards, Washington, D.C., 1976, pp. 269–272. 31. A.M. Gokhale and E.E. Underwood, “A General Method for Estimation of Fracture Surface Roughness: Part I. Theoretical Aspects”, Metallurgical and Materials Transactions A, May 1990, vol. 21A, pp. 1,193–1,199. 32. E.E. Underwood, “Stereological Analysis of Fracture Roughness Parameters”, Acta Stereologica, 1987, vol. 6 Supplement 2, pp. 169–178. 33. A.M. Gokhale and E.E. Underwood, “A New Parametric Roughness Equation for Quantitative Fractography”, Acta Stereologica, 1989, vol. 8, no. 1, pp. 43–52. 34. H.E. Exner and M. Fripan, “Quantitative Assessment of Three-Dimensional Roughness, Anisotropy, and Angular Distributions of Fracture Surfaces by Stereomtery”, Jour nal of Microscopy, 1985, vol. 138, part. 2, pp. 161–178. 35. W.J. Drury and A.M. Gokhale, “Feature Specific Digital Profilometry of Fracture Sur faces”, Quantitative
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2009 4TH INTERNATIONAL BRAZING & SOLDERING CONFERENCE & EXHIBITION April 26–29 Anaheim, CA asmcommunity.asminternational.org/content/Events/ lightmetals09/
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Volume 45, Issue 2, 2009 International Journal of Powder Metallurgy
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