Faster sintering and lower costs with ultra-fine MIM powders Recent advances made in Japan in the production of ultra-fine stainless steel powders mean improvements in MIM processing can be realised…
etal injection moulding or MIM is a segment of the broader field of powder injection moulding (PIM). It is a relatively new technology that uses the shaping advantages of plastic injection moulding but expands the applications to numerous high-performance metals and alloys, as well as metal matrix composites and ceramics [1, 2]. The MIM process consists of mixing a small amount of organic material – the binder phase - with the desired inorganic powder (metals or alloys) to create a feedstock that can flow like plastic under temperature and pressure. This feedstock can be injection moulded into a “green” shape that is an oversized replica of the final part. Generally the organic binder is removed during a step known as debinding, though in some applications the as moulded part is the final component. After debinding, the part is consolidated to high densities which are typically greater than 96 per cent of theoretical density of the metal or alloy. This can be achieved by pressureless sintering (high temperature treatment) or pressure assisted sintering. In this way the MIM process provides designers and engineers with a powerful material shaping technique that can form
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metals and alloys into extremely complex shapes without any metal removal steps such as machining, milling or drilling. Numerous variations of the MIM process are practised by different companies which reflect the different combinations of metal or alloy powders, multi-component organic binders, different moulding techniques and widely diverse debinding processes. The final consolidation step of sintering is generally similar for most MIM applications, with variations being primarily dictated by the material and powder characteristics. The MIM process can be divided into four main steps: feedstock preparation, injection moulding, debinding, and consolidation. The major differences in MIM processing techniques are dictated by the initial choice of the organic binder systems, which in turn dictates the debinding process used to remove the organic binder. The binder systems that are in currently in commercial use are based primarily on wax-polymers, oil-wax-polymer, water-gel, polyacetal and water-polymer. Debinding techniques are usually tailored to ensure clean removal of organic binders, and this has been responsible for the myriad variations of debinding processes - catalytic debinding, pure thermal debind-
ing, wicking, drying, supercritical extraction, organic solvent extraction, waterbased solvent extraction and freeze drying among them. The choice of the debinding equipment is dictated by the choice of the debinding technique used, and it eventually impacts the cost of producing the final part. Once the injection moulded part has been debound (and generally presintered) to ensure that all the organic binder has been removed, the consolidation of the parts is typically carried out in conventional furnaces. Typically, consolidation of most ferrous materials (Fe-Ni alloys, stainless steels, low alloy steels, etc.) and several non-ferrous alloys (nickel and cobalt-based alloys, tungsten alloys, etc.) are sintered in furnaces using some form of reducing atmosphere - typically hydrogen or a mixture of hydrogen with other gases. There are significant differences in the furnaces that are used for consolidation. They include batch furnaces used only for sintering, batch furnaces capable of debinding and sintering, continuous furnaces, and, more latterly, microwave sintering furnaces. Though processing variations are necessary due to the initial choice of the organic binder system, the choice of the
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metal or alloy powder dictates the final properties of the consolidated part. It is quite obvious that the use of different metals or alloys will result in widely different properties in the final component. For example, the use of a 2Ni-98Fe alloy will not have the same corrosion resistance as a stainless steel alloy, while a stainless steel alloy will not have the same strength as a tungsten carbide based material. However, what is not so obvious is the fact that significant property variations can be achieved with the same metal or alloy processed under the same conditions (especially the consolidation conditions), simply by using starting powders with different characteristics. Also, when using different powders, it should be possible to attain similar properties in the same metal or alloy system even when using consolidation conditions that are different. In general, a finer powder when sintered under similar conditions (same temperature, heating rate, and time) will result in a part with higher sintered density, better mechanical properties, and smoother surface finish as compared to a coarse powder of the same alloy. These characteristics of finer powders can be exploited by metal injection moulding to open up new applications and improve the properties of existing applications. Finessing the issue of yields In the past, achieving high superfine powder yields was a major issue which impacted the cost of the finer powders by making it cost-prohibitive. The improvements brought about by ultra-high pressure water atomisation have recently led to the availability of very fine powders for the MIM industry at a reasonable cost. This could provide a major breakthrough in the area of powders for the MIM industry, and it will be one that will have a positive impact on the overall industry. Among the PIM materials that are currently in commercial production, stainless steel is perhaps the most important. Though, Fe-Ni-based alloys were also quite popular in the early days of the PIM industry, with the availability of fine MIM stainless steel powders the volume of stainless steel powders used by the industry increased substantially. Fe-Ni-based alloy used to be based primarily on elemental powder mixes. These required homogenisation along with densification of the material. Incomplete homogenisation
resulted in property variations. This problem has been overcome through the development of fine prealloyed Fe-Ni-based alloys . Within the stainless steel alloy family, the 316L and 17-4 precipitation hardened are the two most popular alloys. The furnaces used for sintering these stainless steels are either batch or continuous with the preferred sintering atmosphere being typically a reducing one. Over the years, there has been little change in the sintering method for MIM stainless steels, although several different atmospheres have been used to sinter 316L steel [4, 5]. The introduction of very fine stainless steel powders could bring about change in his area. A preliminary discussion of the sintered density attained with ultrafine powders was recently reported . This study looked at the processing and properties of two 316L stainless steel powders. The first powder was conventional with a mean particle size of around 10 µm, while the second was ultrafine powder with a mean powder particle size of around 5 µm. The powders used were ultra-high pressure water atomised stainless steel powders from Atmix, Japan. The first powder designated as SUS316L PF-15 was a conventional material with a mean particle size in the range of 7 to 9 µm. The second powder was a superfine powder that had a mean particle size in the range of 3 to 5 µm and was designated as SUS316L PF-5. Particle size measurements on the two powders were performed using the laser diffraction method (Microtrac, Inc., HRA 9320-X100). The tap density of the powder was measured using two different methods. In the first, the height of 100gm of powder taken in a 100 ml cylinder and tapped around 300 times was recorded, but due to surface unevenness it
was impossible to get a correct reading. This was then modified to include an attachment that flattened out the surface of the powder after 300 taps. After flattening, it was followed by another 100 taps. Even with the modification, there was still significant variation. An alternate method was devised to eliminate the effect. This divided the cylinder into two sections and set a slide in the upper section after tapping to smooth the powder surface. This resulted in repeatable results with excellent batch to batch consistency. The Tap Density results reported here were obtained from this method. The specific surface area was measured by the BET method (Mountech Co. Ltd., Macsorb HM model-1201). The theoretical density of the powder was also measured using a gas pycnometer (Micromeritics' AccuPyc Pycnometer). It should be realised that the pycnometer density provides a measurement of the powder density which is generally lower than the density of the metal itself due to the adsorbed moisture and dissolved gases (oxygen and nitrogen). The finer powder is expected to have more gases in solution
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Table 1: Particle size analysis, tap density, and specific surface area of the two powders. D10 (µm)
Tap Density, g/cc
Pycnometer Density, g/cc
Specific Surface Area, m2/g
Atmix SUS316L PF-15
Atmix SUS316L PF-5
Table 2: Chemical composition of the two powders. Sample Designation
Atmix SUS316L PF-15
Atmix SUS316L PF-5
as well as adsorbed gases on the surface due to the higher surface area. Though the powder density was measured to be in the range of 7.85 to 7.89 g/cc, the measured density from the ladle was around 7.95
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g/cc. This latter density was assumed to be the theoretical density of the 316L composition used. The result of the particle size analysis, tap density, pycnometer density, and specific surface area of the two powders are shown in Table 1. The detailed chemical compositions of the powders are given in Table 2. The Scanning Electron photomicrographs of the two powders are shown in Figure 1. Each of the two powders was mixed with a proprietary organic binder to produce the desired feedstock. The mixing was carried out in a kneader for one hour to produce the feedstocks. The moulding of the tensile bars was carried out in an injection moulding machine (Nissei Plastic Industrial). The debinding was carried out in a nitrogen atmos-
phere at 475°C. The total debinding time was around 20 hours. After debinding, the tensile samples were removed for presintering. A pre-sintering step was used to ensure that there was absolutely no binder remaining in the samples. The presintering was carried out in nitrogen using a ramp rate of 5°C/min and a onehour hold. The sintering was carried out at several different sintering temperatures ranging from 900 - 1350°C, using an Argon partial pressure in the range of 100-500 Pa (1-5Torr). The 900°C and 950°C sintering temperatures were used only for sintering the superfine powder, while both the powders were sintered at all the other temperatures of 1000, 1050°C, 1100°C, 1200°C, 1300°C, and 1350°C. A constant hold time of two hours at the maximum sintering temperature was used for all the sintering runs. The sintered densities of the parts were measured by water immersion. The surface roughness of the parts sintered at 1000°C, 1050°C, 1100°C, 1200°C, and 1300°C were measured using a contact type surface roughness measuring instrument (Taylor Hobson). The injection moulded and sintered tensile bars were subjected to tensile testing. For each sintering condition, five tensile bars were pulled to failure using a rate of 3mm/ min. The ultimate tensile strengths and tensile elongations of the sintered tensile bars were determined for several sintering conditions. Some of the as-sintered bars were sectioned for microstructural studies. The sectioned samples were mounted, polished, etched with Aqua Regia and observed in an optical microscope. As the powder size becomes finer, the internal friction of the powder particles is increased. This in turn typically translates to a lower tap density and apparent density (not measured in this case) for the finer powder compared to a coarser powder of similar shape. Also, as the powder shape remains the same but the powder particle size is decreased, the specific surface area of the powder will increase. All of the expected trends were followed by the two powders used in this study as shown by the data in Table 1. As a result of the increased surface area and lower tap density of the finer powder, the viscosity of the material was expected to be higher with the finer powders. However, the nearspherical shape of the powder particles Page 22
Figure 1. a) SEM of SUS316L PF-15; b) SEM of SUS316L PF-5
Figure 2. Relationship between sintering temperature and sintered density of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts.
Figure 3. Relationship between sintering temperature and tensile strength of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts.
formed by the ultra-high pressure water atomisation, results in a lowered viscosity compared to conventional irregular shaped water-atomised powder.
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One of the key advantages of using finer powder is the attainment of higher density at a particular sintering temperature. Figure 2 shows the sintered
density of the metal injection moulded parts sintered at different sintering temperatures for the two different powders. It can be seen that even at extremely low sintering temperatures (for 316L sintering) of 1100°C, the parts using the ultrafine powder (SUS316L PF-5) has already attained a sintered density that is greater than 7.7 g/cc which is almost 97 per cent of theoretical. In contrast, the coarser powder (SUS316L PF-15) sintered at the same temperature exhibited a sintered density of around 7.2 g/cc, which is only around 91 per cent of theoretical. The finer particle size of the ultra fine powder provides a significantly higher sintering potential, which in turn translates into higher sintered density under the same sintering conditions. It can also be seen that in order to attain a density of around 7.7 g/cc with the conventional powder (SUS316L PF-15), the sintering temperature has to be over 1300°C. Many of the metal injection moulded parts had a sintered density requirement of around 7.7 g/cc. The ability of the finer powders to achieve that density at a lower temperature would be a major advantage to part producers, especially for parts that do not have very high strength requirement. Figure 3 shows the relationship between sintering temperature and the ultimate tensile strength of the metal injection moulded parts fabricated from the two powders. The strength initially shows a sharp increase with sintering temperature for the ultra fine powder. A substantial increase in strength is observed when the sintering temperature is increased from 900°C to 1000°C. In fact, the peak tensile strength of the parts made from the ultra fine powder is seen at a sintering temperature of 1050°C, after which the tensile strength shows a slight decrease. In contrast, the coarse powder (SUS316L PF-15) does not show a peak for the tensile strength in the sintering temperature range used for this investigation. The strength curve for the coarse powder samples almost flattens out after a sintering temperature of 1200°C is reached. If one considers the sintering temperature of 1300°C, the difference in the tensile strengths between the ultra fine and the coarse powder samples is almost 50 MPa. However, the difference between the maximum strength achieved Page 24
Figure 4. Relationship between sintering temperature and tensile elongation of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts.
by samples made from the two powders is more than 100 MPa . The variations in the tensile elongation with sintering temperature for the two powders are shown in Figure 4. The tensile elongations, however, did not follow the same trend as the tensile strength properties as shown in Figure 3. The tensile elongation for both the powders is seen to increase with increasing sintering temperature, though there is a general flattening out of the curves at the higher temperatures. The elongation of the parts made from ultra fine powder increases rapidly when the sintering temperature was increased from 950°C to 1000°C. However, even though the tensile strength for the ultra fine powder decreases after 1050°C, the tensile elongation continues to increase up to a temperature of 1200°C after which is flattens out. It can be seen that around 50 per cent of elongation in the ultra-fine samples is attained at a sintering temperature of 1100°C.
Figure 5a: SEM of SUS316L PF-5 sintered at 1000oC.
Figure 5b: SEM of SUS316L PF-5 sintered at 1300oC.
Figure 5c: SEM of SUS316L PF-15 sintered at 1000oC.
Figure 5d: SEM of SUS316L PF-15 sintered at 1300oC.
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Refences 1. Randall M German and Animesh Bose, Injection Molding of Metals and Ceramics, 1997, Metal Powder Industries Federation, Princeton, NJ. 2. Beebhas C Mutsuddy and Renee G Ford, Ceramic Injection Molding, 1995, Chapman & Hall, London, UK. 3. Hisataka Toyoshima, Tokihiro Shimura, Atsushi Watanabe and Hidenori Otsu, “Sintered Compact Properties of Pre-alloyed 2%Ni-Fe Water Atomized Powder” Journal Japan Society of Powder Metallurgy, 2005, vol. 52, no. 6, pp. 437-441. 4. J C Rawers, F Croydon, E A Krabbe, and N W Duttlinger, “Tensile Characteristics of Nitrogen Enhanced PIM 316L Stainless Steel,” Advances in Powder Metallurgy and Particulate Materials, Compiled by M Phillips and J Porter, Metal Powder Industries Federation, Princeton, NJ, 1995, vol. 6, part 6, pp. 229-242. 5. G R White and R M German, “Effect of Process Conditions on the Dimensional Control of Powder Injection Molded 316L Stainless Steel,” Advances in Powder Metallurgy and Particulate Materials, Compiled by C Lall and A J Neupaver, Metal Powder Industries Federation, Princeton, NJ, 1994, vol. 4, pp. 185-196. 6. Hisataka Toyoshima, Minoru Kusunoki, and Isamu Otsuka, “Sintering properties of high-pressure water atomized SUS 316L ultra fine powder,” Proceedings of the PM World Congress, Pusan, Korea, 2007. 7. Materials Standards for Metal Injection Molded Parts, 2007 Edition, MPIF Standards 35, Publisher, MPIF, Princeton, NJ, p.19, 2007. This is in the typical range of MIM properties reported in the MPIF standard . Both the sintered density and the tensile strength of the two powders show a similar variation with sintering temperature. The initial increase in the density and strength is quite rapid at the early stage. The ultra fine powder, due to its associated surface energy, sinters quite rapidly and tends to achieve high density at a much lower temperature. The mass transport and elimination of porosity will be quite rapid due to
the association of the porosity with the grain boundaries. In practice, once the sintering potential was used up and the porosity isolated within the grains, the densification rate quickly levelled off. For the coarser powder, however, the densification rate did not flatten out as dramatically as the ultra fine powder. The drop in the tensile strength in the ultra fine powder MIM parts is likely associated with a rapid grain growth that starts taking place at elevated temperatures. Figure 5a and 5b shows the
microstructures of the ultra fine powder (SUS316L PF-5) MIM parts sintered at 1000°C and 1300°C, respectively. The large difference in the grain size is easily observed from the two microstructures. This grain growth is primarily responsible for the lowering of the tensile strength of the samples at the elevated temperatures. Figure 5c and 5d shows the microstructures of the conventional powder (SUS316L PF-15) MIM parts sintered at 1000°C and 1300°C, respectively. The lower sintering temperature shows the presence of high amount of porosity and prior particle boundaries. The prior particle boundaries were not discernable in case of the ultra fine powder MIM part sintered at 1000°C. At the higher sintering temperature of 1300°C the coarse powder also showed very large grain size. The surface roughness of the parts sintered under the same sintering conditions decreased as the powder particle size became finer. This was an expected trend as the prior particle boundaries will depend on the initial powder particle size and was expected to influence the surface roughness of the part. It can be seen that an increase in the sintering temperature causes a slight increase in the surface roughness of the materials. Figure 6 shows the surface roughness variation with sintering temperature for the parts fabricated from the conventional and ultra fine powders. It can be observed that the use of the ultra-fine powder will result in significantly better surface finish of the final part. It can also be concluded that the use of the ultra-fine powder will result in not only better surface finish of the part but will also result in more complete fill in parts that have significantly finer details.
The Authors THIS article is based on Metal injection molding of ultra-fine 316l stainless steel powders, a paper by Animesh Bose¹, Isamu.Otsuka², Takafumi Yoshida² and Hisataka Toyoshima². ¹ Materials Processing Inc, 5069 MLK Freeway, Fort Worth, TX – 76109, USA ² Epson Atmix Corporation, 4-44 Kaigan, Kawaragi, Hachinohe-shi Aomori-ken, 039-1161 Japan Figure 6. Relationship between sintering temperature and surface roughness of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts.
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