METAL INJECTION MOLDING OF ULTRA-FINE 316L STAINLESS STEEL POWDERS

METAL INJECTION MOLDING OF ULTRA-FINE 316L STAINLESS STEEL POWDERS Animesh Bose, Isamu.Otsuka*, Takafumi Yoshida*, Hisataka Toyoshima* Materials Processing, Inc., 5069 MLK Freeway, Fort Worth, TX 76109, USA * Epson Atmix Corporation, 4-44 Kaigan, Kawaragi, Hachinohe-shi Aomori-ken, 039-1161 Japan ABSTRACT Stainless steels have become one of the most popular materials for the metal injection molding (MIM) process. Due to the poor yield of fine particles in powder production, the cost of relatively fine powders was extremely high. With improvements in the powder processing technique, relatively fine stainless steel powders (mean 10 to 15 micrometers) became available during the 1990 s, resulting in a concomitant decrease in the powder price and a rapid increase in the stainless steel powder consumption in the MIM industry. This situation has remained stable till date. However, the demand for finer and finer powders remained quite strong. This demand for ultra-fine powders for MIM (mean around 5 micrometer or less) was fueled by increasing interest in the area of micro-mim components, the demand for full density parts in the medical field, and the general demand for higher performance stainless steel components. Due to the major improvements (high yields of fine powders) that have been made recently in powder production by Atmix, the use of ultra-fine stainless steel powder has now become a distinct reality. This paper discusses the processing and properties of metal injection molded ultra-fine 316L stainless steel powders. The paper will discuss the advantages in the MIM processing and improvements in the properties of an ultra-fine 316L stainless steel powder (mean around 5 micrometers) compared to a standard stainless steel powder (mean size around 10 micrometers). INTRODUCTION Metal injection molding (MIM) which is a segment of the broader field of powder injection molding (PIM) is a relatively new technology that uses the shaping advantage of plastic injection molding but expands the applications to numerous high performance metals and alloys, as well as metal matrix composites and ceramics [1,2]. The process of MIM consists of mixing a small amount of organic material (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 molded 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 molded part is the final component. After debinding, the part is consolidated to high densities (typically greater than 96% of theoretical density of the metal or alloy) by pressureless sintering (high temperature treatment) or pressure assisted sintering. Thus, the MIM process provides the designers and engineers with a powerful material shaping technique that can 1

shape metals and alloys into extremely complex shapes without any metal removal steps such as machining, milling, drilling, etc. Numerous variations of the MIM process is practiced by different companies which is truly a reflection of the different combinations of metal or alloy powders, multi-component organic binders, different molding techniques, 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 molding, debinding, and consolidation. The major differences in the MIM processing techniques are dictated by the initial choice of the organic binder systems, which in turn dictates the debinding process that is 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, water-polymer, etc. The debinding techniques are often tailored to ensure the clean removal of the organic binders, which is responsible for the variations in the debinding processes (catalytic debinding, pure thermal debinding, wicking, drying, supercritical extraction, organic solvent extraction, water-based solvent extraction, freeze drying, etc.). 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 molded part has been debound (and generally presintered) to ensure that all the organic binder has been removed, the consolidation of the parts have been 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.) has been sintered in conventional 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 and they include batch furnaces used only for sintering, batch furnaces capable of debinding and sintering, continuous furnaces, and even microwave sintering furnaces. Though processing variations are necessary due to the initial choice of the organic binder system, the choice of the 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 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 different starting powder characteristics. Also, when using different starting powder characteristics, 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 molding to open up new applications and improve the properties of existing applications. In the past, achieving the 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 2

water atomization 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 one. Though, Fe-Ni-based alloys have also been 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. The Fe-Ni-based alloy, in the past was primarily based on elemental powder mixes. This required the homogenization along with densification of the material. Incomplete homogenization resulted in property variations. This problem has been overcome through the development of fine prealloyed Fe-Nibased alloys [3]. Within the stainless steel alloy family, the 316L and 17-4 precipitation hardened are the two most popular alloys. The furnaces used for sintering this 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 that has been used for the sintering of metal injection molded stainless steels. There have been several sintering atmospheres used to sinter MIM 316L stainless steel parts. Some of the sintering atmospheres include the use of partial pressure of hydrogen, argon, or nitrogen [4]. The use of a mixture of hydrogen and nitrogen or even pure flowing hydrogen has also been used [5]. The introduction of the very fine stainless steel powders could also bring about some changes to the sintering conditions that have typically been used by the industry. A preliminary discussion of the sintered density attained with the ultrafine powders was recently reported [6]. This paper will discuss the processing and properties of two 316L stainless steel powders. The first powder will be the conventional one with a mean particle size of around 10 µm, while the second will be the ultrafine powder with a mean powder particle size of around 5 µm. EXPERIMENTAL Two 316L stainless steel powders were used for this experiment. The powders used were ultrahigh pressure water atomized 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 method, the height of 100 gm of powder taken in a 100 ml cylinder and tapped around 300 times is recorded. In this case, due to the 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 and after flattening, it is followed by another 100 taps. Even with the modification, there was still significant variation. An alternate method was devised to eliminate the effect of uneven surface. This process divides the cylinder into two sections and sets 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 are obtained from this method. The specific surface area was measured by 3

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 realized 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 the dissolved gases (oxygen, nitrogen). The finer powder is expected to have more gases in solution as well as adsorbed gases on the surface due to the higher surface area. Though the powder density is measured to be in the range of 7.85 to 7.89 g/cc, the measured density from the ladle is around 7.95 g/cc. This latter density is assumed to be the theoretical density of the 316L composition that is 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. Table 1: Particle size analysis, tap density, and specific surface area of the two powders. Sample Designation D10 (µm) D50 (µm) D90 (µm) Atmix SUS316L PF-15 Atmix SUS316L PF-5 Tap Density, g/cc Pycnometer Density, g/cc Specific Surface Area, m 2 /g 3.2 8.5 19.1 4.4 7.89 0.27 2.1 4.0 7.3 3.9 7.85 0.47 Table 2: Chemical composition of the two powders Sample Designation Cr wt.% Ni wt.% C wt.% Fe wt.% Atmix SUS316L PF-15 Atmix SUS316L PF-5 O wt.% N wt.% 16.50 12.53 0.026 Bal. 0.39 0.07 16.66 12.49 0.021 Bal 0.40 0.04 a 10 μm b Figure 1. a) SEM of SUS316L PF-15; b) SEM of SUS316L PF-5 4

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 1 hour to produce the feedstocks. The molding of the tensile bars was carried out in an injection molding machine (Nissei Plastic Industrial). The debinding was carried out in a nitrogen atmosphere using a temperature of 475 o 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 remained in the samples. The presintering was carried out in nitrogen using a ramp rate of 5 degrees centigrade per minute and a 1-hour hold. The sintering was carried out at several different sintering temperatures ranging from 900 to 1350 o C, using an Argon partial pressure in the range of 100-500 Pa (1-5Torr). The 900 and 950 o C sintering temperature was used only for the sintering of the superfine powder, while both the powders were sintered at all the other temperatures of 1000, 1050, 1100, 1200, 1300, and 1350 o C. A constant hold time of 2 hours at the maximum sintering temperature was used for all the sintering runs. The sintered densities of the parts were measured by water immersion technique. The surface roughness of the parts sintered at 1000, 1050, 1100, 1200, and 1300 o C are measured using a contact type surface roughness measuring instrument (Taylor Hobson). The injection molded and sintered tensile bars were subjected to tensile testing. For each sintering condition, 5 tensile bars were pulled to failure. The tensile bars were pulled to failure using a rate of 3 mm/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. Appropriate photomicrographs were taken and have been used in the discussion. RESULTS AND DISCUSSION As the powders 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 are followed in 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 is expected to be higher with the finer powders. However, the near spherical shape of the powder particles formed by the ultra-high pressure water atomization, results in a lowered viscosity compared to conventional irregular shaped water atomized powder. One of the key advantages of the use of finer powder is the attainment of higher density at a particular sintering temperature. Figure 2 shows the sintered density of the metal injection molded parts sintered at different sintering temperatures for the two different powders. It can be observed from the figure that even at extremely low sintering temperatures (for 316L sintering) of 1100 o C, the parts using the ultra-fine powder (SUS316L PF-5) has already attained a sintered density that is greater than 7.7 g/cc which is almost equal to 97% of theoretical density. In contrast, the coarser powder (SUS316L PF-15) sintered at the same temperature exhibits a 5

sintered density of around 7.2 g/cc, which is only around 91% of theoretical density. 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 o C (which is around 200 o C higher than that needed by the ultra fine powder). Many of the metal injection molded parts have a sintered density requirement of around 7.7 g/cc. It can be observed that this density is reached in case of the ultra fine powder even at sintering temperatures of around 1100 o C. This will be a major advantage to part producers, especially for parts that do not have very high strength requirement. 8.0 7.5 Density (g/cc) 7.0 6.5 6.0 SUS316L PF-5 SUS316L PF-15 5.5 5.0 800 900 1000 1100 1200 1300 1400 Sintering Temperature ( C) 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 shows the relationship between sintering temperature and the ultimate tensile strength of the metal injection molded 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 to 1000 o C. In fact, the peak tensile strength of the parts made from the ultra fine powder is seen at a sintering temperature of 1050 o 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 that has been used for this investigation. The strength curve for the coarse powder samples almost flattens out after a sintering temperature of 1200 o C is reached. If one considers the sintering temperature of 1300 o C, the difference in the tensile strengths 6

between the ultra fine and the coarse powder samples is almost 50 MPa. However, if the difference between the maximum strength achieved by samples made from the two powders is concerned, the difference is over 100 MPa. Also, the maximum tensile strength (623 MPa) is achieved in case of the ultra fine powder at a sintering temperature that is significantly lower than the sintering temperature needed for the conventional powder to achieve its highest strength (1050 o C for ultra fine versus 1350 o C for the conventional powder). It should also be pointed out that the peak tensile strength reached by the ultra fine powder samples is around 100 MPa higher than the typical 316L tensile strength value reported in the MPIF Standard [7]. 700 600 Tensile Strength (MPa) 500 400 300 200 100 SUS316L PF-5 SUS316L PF-15 0 800 900 1000 1100 1200 1300 1400 Sintering Temperature ( C) Figure 3. Relationship between sintering temperature and tensile strength of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts. The variations in the tensile elongation with sintering temperature for the two powders are shown in Figure 4. The tensile elongation, however, do 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 is increased from 950 to 1000 o C. However, even though the tensile strength for the ultra fine powder decreases after 1050 o C, the tensile elongation continues to increase till a temperature of 1200 o C after which is flattens out. It should be realized that elongation values around 50% that is attained at a sintering temperature of 1100 o C for the ultrafine powder is in the typical range of MIM properties reported in the MPIF standard [7]. 7

80 70 Elongation (%) 60 50 40 30 20 SUS316L PF-5 SUS316L PF-15 10 0 800 900 1000 1100 1200 1300 1400 Sintering Temperature ( C) Figure 4. Relationship between sintering temperature and tensile elongation of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts. 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 will sinter quite rapidly and would tend 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. Once the sintering potential is used up and the porosity is isolated within the grains, the densification rate is expected to quickly level off. For the coarser powder, however, the densification rate does 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 o C and 1300 o 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 o C and 1300 o 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 o C. At the higher sintering temperature of 1300 o C the coarse powder also showed very large grain size. 8

Figure 5a: SEM of SUS316L PF-5 sintered at 1000 o C. 20μm Figure 5b: SEM of SUS316L PF-5 sintered at 1300 o C. 9

Figure 5c: SEM of SUS316L PF-15 sintered at 1000 o C. 20μm Figure 5d: SEM of SUS316L PF-15 sintered at 1300 o C. 10

The surface roughness of the parts sintered under the same sintering conditions is decreased as the powder particle size becomes finer. This is an expected trend as the prior particle boundaries will depend on the initial powder particle size and is expected to influence the surface roughness of the part. It is also observed 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 two powders, 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. 1.0 0.8 Ra Roughness (µm) 0.6 0.4 0.2 SUS316L PF-5 SUS316L PF-15 0.0 800 900 1000 1100 1200 1300 1400 Sintering Temperature ( C) Figure 6. Relationship between sintering temperature and surface roughness of the ultra fine (SUS316L PF-5) and conventional (SUS316L PF-15) powder MIM parts. CONCLUSIONS This investigation discusses the processing and some properties of two metal injection molded 316L stainless steel powders, ultra fine powder (SUS316L PF-5) and a conventional powder (SUS316L PF-15). The ultra fine powder showed a significantly higher sintering rate and could be sintered at temperatures in the range of 1050 to 1100 o C to achieve high strengths and sintered densities. The surface roughness of parts made from the ultra fine powder was lower than that of the parts made from conventional powder. The peak tensile strength of parts made from the ultra 11

fine powder was achieved at a temperature of 1050 o C, after which the strength decreased most likely due to grain growth. The use of the ultra fine powder will allow the processing of MIM parts at a lower temperature and the parts can be fabricated to achieve higher tensile strength and better surface finish compared to parts made from conventional powders. The ultra fine powder will also allow filling of the fine details in the mold. ACKNOWLEDGEMENT The authors would like to acknowledge the help of Mr. Pinaki Bose in preparing the figures used in this paper and Mrs. Prarthana Bose for her help in editing of the manuscript. REFERENCES 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 highpressure 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. 12