Powder Injection Moulding International December 200966 Vol. 3 No. 4

Technical Paper

IntroductionLow carbon austenitic chromium-nickel-iron alloys with 17-19% chromium and 8-14% nickel are widespread due to their mechanical and technological advantages [1]. 316L with 16.5 – 18.5% chromium and 12-14% nickel is one of the most common materials for MIM-production. This is due to its excellent corrosion resistance, its high ductility, its high toughness and impact strength, even at lower temperatures and its good weldability [1]. Furthermore the material is characterised by its low yield point and its suitability to cold work hardening [1].

The metallurgy of chromium-nickel-iron alloys is rather complicated. The materials are considered corrosion resistant at chromium levels above 12-14% [1,2]. This is due to the passivation property of chromium. A thin protective surface film forms instantly under oxidising atmospheres. The stainless steel is passivated because an oxide film seals the bulk metal from oxygen diffusion and further corrosion [3]. At higher chromium concentrations the corrosion resistance increases and at lower concentrations it decreases. Therefore, the chromium has to be evenly distributed throughout the metal lattice. Total as well as local reduction of chromium will reduce the corrosion resistance.

If powder mixtures of master alloys and carbonyl iron are used alloying appears during sintering. In such powder concepts fine carbonyl iron powders are mixed with so called master alloys. The usually gas atomised master alloy is enriched with chromium and nickel atoms. During sintering the chromium and nickel atoms diffuse into the iron lattice. The chromium and nickel atoms take the atomic place of the iron atoms. This diffusion process takes place at sintering temperature and requires time. If the alloying process is stopped prematurely or if the master alloy carbonyl iron powder mixture is inhomogeneous in the first place the alloy will suffer liquation. In that case the chromium concentration may locally not reach 16.5%. It may even stay below 12% and be subject to

corrosion. For the present study this effect is not relevant because gas atomised powders were used. Each powder particle consists of pre-alloyed 316L. Thus there is neither alloying nor liquation taking place.

Characteristic of 316L is a low carbon content which should not exceed 0.03% [4]. At this low carbon content 316L consists of a non magnetic austenitic face-centered cubic crystal structure. If the carbon content increases the austenitic temperature increases dramatically. At low cooling rates or during annealing, carbon

Influence of sintering conditions and furnace technology on chemical and mechanical properties of injection moulded 316L

In industrial practice, stainless steel injection moulded parts are sintered under various sintering conditions. Common are molybdenum or graphite vacuum furnaces. Different sintering atmospheres and pressures are used. Based on a water soluble binder system 316L test bars are injection moulded and sintered under various sintering conditions (vacuum, nitrogen, hydrogen and argon). The main task of the study is to elaborate the impact of sintering atmospheres on the properties of Stainless Steel. To this purpose metallurgical, chemical and mechanical tests are carried out and discussed. Sintering tests are carried out in graphite and molybdenum furnaces. Debinding and sintering in a furnace with a retort as well as thermal binder removal and sintering in two different furnaces are evaluated. As shown by the tests carbon control during debinding and control of the sintering athmosphere has an important influence on finished part quality. So does the furnace technology used.

Steffen Krug1, Stefan Zachmann2

1 Polymer-Chemie GmbH, Metal Powder Compound Division, Am Gefach, 55566 Bad Sobernheim, Germany2 Nabertherm GmbH, Bahnhofstr. 20, 28865 Lilienthal, Germany

Fig. 1 Nabertherm VHT-08 sintering furnace

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segregates and carbides are created. This effect is rather complicated since the carbide creation may be inhibited by fast cooling. In this case the carbon stays dissolved as an interstitial atom in meta-stable conditions. This effect is of technical relevance since carbide segregations are subject to corrosion and influence the mechanical properties of the alloy [1]. The low temperature segregation appears along the grain boundary. According to theory the carbides are enriched in chromium thereby reducing the chromium level next to the grain boundary. The alloy is losing its corrosion resistance along the grain boundary resulting in inter-crystalline corrosion. However, if the total carbon concentration is below 0.05% there is no such corrosion effect [1].

A further effect of local chromium reduction and hence reduced corrosion resistance is reported when vacuum sintering Cr-Ni-steels [5]. It is particularly Mn, Cu and Cr which evaporate from steels during vacuum sintering. Up to one third of the chromium content may evaporate in the surface area of austenitic Cr-Ni-steels. Operating the furnace at nitrogen or argon partial pressure is reported to reduce the surface chromium evaporation [6].

In other studies [7] vacuum sintering is reported to provide the most corrosion resistant material since there is no chromium depletion through the formation of chromium oxides or nitrides. This is due to the low stability of chromium oxide under reduced pressure at temperatures exceeding 900°C. The effect is known as vacuum reduction [6]. The oxygen content is reported to decrease with increasing hydrogen in the sintering atmosphere and with higher sintering temperatures [8].

In an early study from 1955, a pure hydrogen atmosphere during sintering is reported to avoid corrosion problems and to increase the mechanical properties of Stainless Steel [9]. If dissociated ammonia is used as a sintering atmosphere both a low dew point and rapid cooling are required.

Using both nitrogen sintering atmospheres and slow cooling below 1000°C was reported to create chromium nitride in the near surface region, leading to a significant reduction in corrosion resistance and a slight improved crushing strength. The amount of chromium nitride formed is dependent on the bulk nitrogen content. A critical absorbed nitrogen content was in excess of 0.4%. A higher sintering temperature, faster cooling rates and lower atmosphere nitrogen content are all reported to reduce absorbed nitrogen [8]. At cooling rates faster than 200 K/min chromium nitride is claimed to be totally avoided [7].

Ignoring the influence on corrosion resistance, sintering 316L in a nitrogen gas environment is reported to increase yield and tensile strength with increasing nitrogen concentration in the process gas. Yield strength increased from 170 MPa at 0.006% nitrogen to 388 MPa at 0.31% nitrogen [10].

One important finding was that corrosion resistance and mechanical behavior varied remarkably between different powder vendors [11].

In general the corrosion resistance and mechanical properties of sintered austenitic stainless PM steel are reported to be lower than its corresponding cast steels [8,12,13]. Besides this, the increase of carbon, nitrogen and oxygen levels as well as the remaining porosity will influence the mechanical properties and reduce corrosion resistance.

Experimental detailsThe feedstock used in this study is commercially available and distributed as polyMIM 316L D 120E by Polymer-Chemie GmbH in Germany. The 316L powder is pre-alloyed and gas-atomised with a grain size of d80 > 22µm. The chemical composition of the raw powder is given in Table 1.

Small tensile test bars (Fig. 3) are injection moulded using a commercial injection moulding machine. The mouldings are water binder removed at 40°C for 10 h reaching a weight loss of 4.3%. After water binder removal the samples are dried in an air circulating furnace at 80°C for 4 hours. A comprehensive description of water binder removal from polyMIM mouldings was published in an earlier study [14].

For sintering a Nabertherm VHT-8 Graphite and a VHT-8 Molybdenum furnace (Fig. 1) were used. Both furnaces are equipped with a retort. During thermal binder removal the chamber was purged at 150 l/h. The process gas was hydrogen, nitrogen or argon. After a dwelling time of 2 hours for thermal binder removal at 600°C the process gas was purged into the retort. In order to have the same sintering profiles in both furnaces and for the different sintering atmospheres the temperatures are taken from a sample thermocouple inside the retort. The set sintering temperatures are adjusted to reach a sintering temperature of 1350°C for the samples in all experiments. Holding time was 2 hours for all sintering experiments. Control of gas flow during thermal debinding and sintering is described later (see section Residual binder and furnace technology).

For each sintering run five tensile test bars were thermally binder removed and sintered in nitrogen, hydrogen and argon. In the

Chemical Composition

Fe Cr Ni Mo Mn O N C

bal. 16.7 10.6 2.5 1.3 0.131 0.205 0.015

Table 1 Chemical composition of the 316L powder

Fig. 2 Stress-strain behavior of 316L sintered in a molybdenum furnace at different sintering atmospheres

Fig. 3 In nitrogen and hydrogen sintered 316L test bars before and after pulling

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graphite furnace additional samples were sintered using N2/H2 95/5 under vacuum. For vacuum sintering the samples were thermally binder removed at an atmospheric nitrogen pressure and sintered without gas flow at about 1x10-2 mbar.

For mechanical testing a Karl Frank type 81801D was used. Oxygen and nitrogen concentration were measured using a Leco TC400. The carbon concentration was measured using a Leco CS230.

Results and discussionSintering in a molybdenum furnaceFig. 2 shows the stress-strain behaviour of 316L sintered at 1350°C in a molybdenum furnace when using hydrogen, nitrogen or argon. Fig. 3 shows the corresponding test bars before and after pulling. The hydrogen and argon samples show identical mechanical behaviour. As expected the 316L experiences a gradual elastic-plastic transition having a yield strength of around 180 MPa and a maximum strength of 560 MPa at an elongation of about 65%. The mechanical values exceed the material characteristics given in MPIF MIM standards [4].

The 316L sample sintered in nitrogen shows a different mechanical behaviour. The yield point clearly increases to about 330 MPa and appears to be more pronounced. The elongation is reduced to 45% and the neck formation is almost eliminated. Comparing the samples, the nitrogen sample clearly experienced a strengthening during sintering.

As expected for 316L the argon and hydrogen samples show a moderate amount of necking. This plastic deformation corresponds to the motion of large numbers of dislocations. The 316L samples sintered in nitrogen, however, show reduced elongation which indicates that the lattice hosts some impurity atoms.

Looking at the interstitial element contents (Table 2) oxygen and carbon levels are low and similar for all sintering atmospheres. However, there is clear evidence of nitrogen absorption from the sintering atmosphere if 316 L is sintered using nitrogen as a process gas. The nitrogen content increased from 0.2% to 0.38% during sintering in proportion to the nitrogen content of the raw powder. Looking at the nitrogen content measured from argon and hydrogen samples the nitrogen concentration is remarkably reduced to 0.002% to 0.001% respectively. These findings clearly show both nitrogen absorption and desorption are possible during sintering of stainless steel. At low nitrogen partial pressure in the sintering atmosphere, nitrogen is desorbed from the mouldings. At increasing nitrogen partial pressure nitrogen is absorbed.

As reported elsewhere [10] the mechanical properties are improved by increasing the nitrogen concentration. All samples show a shiny metal surface. However, samples sintered in nitrogen are somehow darker than the other samples indicating a different surface chemistry. Within the scope of this study the influence on corrosion resistance remains open and will be investigated in subsequent studies.

Sintering Results

atmospheredensity (g/cm3)

yield strength

(MPa)

tensile strength

(MPa)

fracture strain %

N % O % C %

Moly. Furnace

hydrogen 7,92 178 554,78 81,4 0,001 0,015 0,0038

argon 7,75 181 543,33 78,4 0,002 0,02 0,0037

nitrogen 7,13 330 615,16 44,9 0,383 0,05 0,0059

Graphite Furnace

argon 7,74 182 554,46 76,8 0,002 0,06 0,0061

nitrogen 7,22 375 680,09 53,3 0,381 0,06 0,0036

N2/H2 (95/5) 7,25 350 640,75 31,2 0,367 0,02 0,0864

argon (pre-debinded)

7,68 182 548,3 69,9 0,002 0,05 0,0036

nitrogen (pre-debinded)

7,05 375 633,21 39,25 0,381 0,05 0,0035

vacuum (1x10-2 mbar)

7,92 190 548,81 81,2 0,001 0,02

Fig. 4 Optical micrographs of 316L sintered in hydrogen (left), argon (centre) or nitrogen (right) after etching

Table 2 Mechanical and chemical results of sintered 316L samples at various sintering atmospheres

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Fig. 4 shows optical micrographs of 316L samples sintered in a molybdenum furnace. All samples show an austenitic microstructure at comparable grain sizes. All samples are sintered at the same thermal profile. The sample sintered in nitrogen is clearly distinguished by a higher porosity. The hydrogen samples reached a density of 7.92 g/cm3, the argon samples reached a density of 7.75 g/cm3 and the nitrogen samples reached only 7.13 g/cm3. The small density variation between sintering in argon or hydrogen may explain the small variations in mechanical behavior evident in Fig. 2 and Table 2. The elimination of porosity in the final stage of sintering is inhibited because argon remains in the final pores and builds an internal pressure. Other than the argon atoms the small hydrogen atoms are more likely to diffuse into the metal lattice and therefore do not inhibit the elimination of final porosity.

For the nitrogen samples the sintering activity was clearly impeded, presumably by the desorbed nitrogen which obviously reduces the mass transport mechanism during sintering. By increasing the sintering temperature from 1350°C to 1365°C the density of the nitrogen-sintered sample could be increased from 7.13 g/cm3 to 7.45 g/cm3. However, the density of vacuum or hydrogen sintered samples could not be reached.

Any evidence of Cr2N could not be established witin this study. At 0.38% the total nitrogen content was just below the critical content of 0.4% [8]. At the given cooling rate a Cr2N creation would be possible because cooling rates are usually not higher than 15-25 K/min for such furnaces. The fast cooling which is reported to prevent Cr2N creation at cooling rates above 200 K/min [7] is technologically not feasible for molybdenum sintering furnaces.

Sintering in a graphite furnaceFig. 5 shows the stress-strain behavior of 316L sintered at 1350°C in a graphite furnace using nitrogen, argon, vacuum and a gas mixture of 5% hydrogen and 95% Nitrogen (N2/H2 95/5). The argon and vacuum sintered samples show the same mechanical behaviour and are comparable with the samples sintered in the molybdenum furnace. Looking at the nitrogen, oxygen and carbon content they are also comparable to the results obtained from the molybdenum samples.

As expected the samples sintered in nitrogen show the same results as with sintering in a molybdenum furnace. The increased nitrogen content of 0.38% causes an increase in yield strength and a decreasing elongation. The correlation is considered to be the same as in molybdenum furnaces described above.

Using a nitrogen-hydrogen mixture in the graphite furnace, however, seems to yield a different result from the hydrogen samples in the molybdenum furnace. Whilst the oxygen content is comparably low, the carbon level increased from about 0.015% (powder) to 0.086% in the sintered samples. The nitrogen content increased to the same level as sintering in pure nitrogen. As a result the yield strength increases to 350 MPa and the elongation decreases to 31%.

For both furnaces the carbon level was reduced from 0.015% in the starting powder to about 0.0038% to 0.0059%. The reduction of carbon during sintering is due to a combined reduction of oxides and carbides which are forming carbon monoxide escaping into the sintering atmosphere. In the molybdenum furnace a hydrogen atmosphere results in a better carbon reduction. The intention of using a nitrogen-hydrogen mixture at 5% hydrogen in the graphite furnace was to provide an improved reducing sintering atmosphere for the 316L samples. However, it causes an increase in the carbon level. A relation to the residual binder from the thermal debinding step can be excluded. This is based on the fact that 316 L samples free of binder experience the same carbon increase in graphite furnaces as if sintered in a nitrogen-hydrogen atmosphere. Neither

of the other sintering experiments, irrespective if sintering takes place in a molybdenum or a graphite furnace, showed such a high carbon level.

The use of hydrogen in a graphite furnace above 900°C is known to create a partial pressure of methane [15]. Furnaces made from graphite fibre board insulation are known to absorb moisture and gathering water on the fibre surface [15]. During sintering the fibre insulation desorbs water to the sintering atmosphere. Methane is known to react with water and form carbon monoxide [6,16] which causes a shift in carbon level during sintering. Certainly, the extent of this study does not prove the relevance of this assumption. However, the use of atmospheres containing hydrogen in graphite furnaces certainly caused a carbon increase in the sintered 316L samples.

Micrographs of the 316L samples sintered in the graphite furnace do not show any differences to the micrographs of samples sintered in a molybdenum furnace as shown in Fig. 3. Of particular interest is the high density of the vacuum sintered samples. Sintering under vacuum in a graphite furnace yielded almost the same density as sintering in a molybdenum furnace. Using argon or vacuum in a graphite furnace results in having the same mechanical and chemical properties of 316L as when sintering in molybdenum furnaces. The impact on corrosion resistance will be the subject of further studies.

Residual binder and furnace technologyLooking at the carbon level there was no evidence of an increased carbon level in the sintered samples due to residual binders. It shows

Fig. 5 Stress-strain behavior of 316L sintered in a VHT graphite furnace using nitrogen, argon, vacuum and an argon/hydrogen in comparison

Fig. 6 Schematic drawing of the sintering furnace

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that there is no carbon increase due to thermal decomposition of the used binder system. This fact was also evident in an earlier study on titanium and titanium alloys using the same binder system [17].

The carbon control during sintering is determined by the furnace technology. If the binders are starting to thermally decompose, which appears to be the case at temperatures starting from 280°C to about 420°C, the furnace is likely to acquire organic contamination. In that case binder residuals may condensate on cold surfaces in the furnace. This may happen in the cold chamber, in the gas outlet or even on the insulation material. If during sintering the temperatures rise these areas may heat up and the binder residuals are likely to evaporate into the sintering atmosphere. The organic contamination in the sintering atmosphere leads to carburisation at higher temperatures during sintering.

Fig. 6 shows a schematic drawing of the used sintering furnace VHT08. During binder removal at temperatures of usually up to 600°C the process gas is purged into the chamber outside the retort and the gas outlet inside the retort is opened. During this thermal debinding step all binders are transported from the retort via a heated gas outlet directly into an after burner. At this stage atmospheric pressures are required for a better heat transport mechanism in the furnace and in order to minimise counter diffusion of binders from the retort into the chamber. Using this gas flow direction the furnace stays clean of contamination during thermal binder removal.

In order to reduce any contamination inside the retort during sintering the gas flow will be adjusted. Clean and dry process gas is purged into the retort during sintering and escapes from the retort into the heating chamber and from there into the gas outlet. The samples are well-housed in a protective gas atmosphere at higher temperatures. It has been proven in practice that, using this sintering technology, increasing carbon levels do not become an issue in MIM production.

ConclusionsThere is no evidence of the influence of the furnace insulation material in both graphite and molybdenum refectories on the chemical or mechanical quality of sintered 316L MIM components.

Using hydrogen or vacuum during sintering 316L results in high density and ductility. Using argon results in a slight decrease of density.

Using nitrogen causes a nitrogen absorption in the 316L samples. The yield strength increases, and the elongation and the density decreases.

Whilst a nitrogen atmosphere causes a nitrogen absorption, vacuum, argon and hydrogen atmospheres cause a reduction of the nitrogen content in sintered 316L.

The carbon level can be reduced in both furnaces. In the molybdenum furnace hydrogen atmospheres slightly enhance the carbon reduction. In graphite furnaces hydrogen partial pressure causes an unfortunate increase in carbon content in the sintered mouldings. The dissolved carbon leads to an increase of the yield strength combined with an embrittled 316L.

Using argon or vacuum in a graphite furnace results in the same mechanical and chemical properties of 316L as the sintering in molybdenum furnaces.

This study shows that there is no evidence of quality differences in 316L samples sintered in a molybdenum furnace using hydrogen and samples sintered in a graphite furnace under vacuum. Considering that molybdenum hydrogen furnaces are almost double the price compared to a graphite vacuum furnace, this finding has an important impact on the production costs of 316 L components simply due to the different furnace technologies available.

References[1] “Metallographie“ 13 Ed., H. Schumann, Deutscher Verlag für Grundstoffindustrie Stuttgart, 1990.[2] “Werkstofftechnik“ 3 Ed., W. Bergmann, Hanser, 2002.[3] “Materials Science and Engineering an Introduction“ 4 Ed., W. D. Callister, John Wiley& Sons, Inc. 1997.[4] MPIF Standard 35, Metal Powder Industries Federation, Edition 2007.[5] “Processing of sintered stainless steel parts”, Powder Metall. vol. 34 pp.259-261, 1991.[6] “Control of Protective Atmospheres during Sintering”, P. Beiss, Proceedings Euro PM2000, pp. 147-157.[7] “The Corrosion Resistance of Sintered Austenitic Stainless Steel”, R.L. Sands, G.F. Bidmead and D.A. Oliver, Modern Developments in Powder Metallurgy, Vol.2, H.H. Hauserner, Plenum Press, New York, pp 73-83, 1966.[8] “Influence of Sintering Variables on the Corrosion Resistance of 316L Stainless Steel”, G. Lei, R. M. German and H. S. Nayar, Powder Metallurgy International, vol.15, no.2, 1983.[9] “An Evaluation of Dissociated Ammonia and Hydrogen Atmospheres of Sintering Stainless Steel“, H.S. Kalish and E.N. Mazza, Trans. TMS-AIME Vol.203, pp.304-310, 1955.[10] “Nitrogen Enhanced Stainless Steel by Powder Injection Molding”, J. Rawers, F. Croydon, R. Krabbe and N, Duttlinger, Engineering and Technology, Vol.32 No. 4, pp 319-321. 1996.[11] “Press und Sinterverhalten von Sonderstahlpulvern”, G. Jangg, O. Knotek and K. Faber, Planseeber. Pulvermet. vol.16, pp. 194-203, 1968.[12] “Corrosion resistance of alloyed powder metallurgy austenitic stainless steels in acid solutions”, F. Velasco, J.M. Ruiz-Roman, J.M. Torralba, J.M. Ruiz-Prieto, British Corrosion Journal, vol.31, no.4 pp. 296-299, 1996,[13] “Study of the influence of nitric acid concentration on the corrosion resistance of sintered austenitic stainless steel”, E. Otero, A. Pardo, E. Saenz, M. V. Utrilla and P. Hierro, Corrosion Science, vol.38, No.9, pp. 1485-1493, 1996.[14] “Reaction and Transport Kinetics of a Water-Soluble MIM Binder System ”, S. Krug, Advances in Powder Metallurgy and Particulate Materials, MPIF World Congress Washington 2008.[15] “Untersuchungen zur Beständigkeit von Graphitwerkstoffen in Elektroöfen bei Gasdrücken bis 1 bar” V.Uhlig, Technische Universität Bergakademie Freiberg, Shaker Verlag p117, 2000. [16] “Sintering Theory and Practice“, R. M. German, John Wiley & Sons Inc., p.424, 1996. [17] “Titan Injection Molding “, S. Krug, Presentation, PIM 2009 in Orlando.

Technical Paper