The Role of Oxygen Transfer in Sintering of Low Alloy Steel Powder Compacts: A Review of the “Internal Getter” Effect
- 851 Downloads
The chemical aspects of sintering have to be considered, in particular the role of oxygen. For sintered alloy steels used for highly stressed components, traditional alloy elements have been Cu, Ni and Mo, which in their oxygen affinity are very similar to the base constituent iron. Advanced alloying systems however contain Cr, Mn and/or Si. In the present study it is shown that one of the principal aspects of sintering to be considered is oxygen transfer from the base iron oxides to the alloy elements, which then form oxides that are more difficult to reduce. This process, defined as “internal gettering”, occurs both in mixed powder compacts and in prealloyed materials, although through different mechanisms. The effect can at least be alleviated by presintering in H2 in the 400°C range, part of the oxygen being removed as H2O before internal gettering becomes kinetically effective. However, in industrial practice, this collides with delubricaton. Furthermore for both alloy variants high temperature sintering is advantageous because it enhances reduction of the more stable oxides, thus eliminating the effects of internal gettering.
Traditionally, studying the fundamental mechanisms of sintering has been focused on physical processes such as: solid state sintering, densification and grain growth,1 the role of contact geometry,2, 3, 4 of defects5,6 and of liquid phase.1,7 The chemical side, in particular the role of oxygen, has received much less attention by science, except the effect of oxide layers on wetting8 and atmosphere effects in some cases (See Ref. 9). It has been mostly left to the industry, one major task being to take care of removing the oxygen introduced through the powders. Oxides may strongly inhibit formation of stable metallic sintering bridges, as quite dramatically shown when sintering aluminium compacts.10,11 Not surprisingly, the first thorough and systematic studies on the effects of oxygen, in particular its interaction with carbon and the ensuing gas-forming reactions, came from the hardmetal side, performed by Leitner et al.12 With WC-Co, precise adjustment of the carbon content within a very narrow interval is of decisive importance to avoid either formation of eta carbides or of free graphite, both of which render the hardmetal unsuitable.13 Carbon control is therefore an essential precondition for obtaining marketable products in the hardmetals industry.
A special case is the sintering of titanium and titanium alloys, where the surface oxide films on titanium powder particles will actively be dissolved. Oxygen diffuses into the bulk titanium matrix during heating, at temperatures above 700°C, and the oxides disappear before the isothermal sintering temperature is reached.14 This surface “cleaning” occurs however at the expense of ductility. Because of this, oxygen getters or scavengers are often introduced to ensure good ductility of the as-sintered titanium alloys parts.15,16
This changed when other alloy elements with higher oxygen affinity were introduced. The driving force was in part economy, elements such as Cr or Mn being markedly cheaper than Ni and also are priced more stable.19 Furthermore, Ni is regarded toxic, while Cu offers problems during recycling. Therefore, numerous approaches have been made to replace Ni-Cu and Mo by Cr and Mn, both the mixing—in part through masteralloys—and the prealloying routes being studied.20, 21, 22, 23, 24, 25, 26, 27, 28, 29 However, processing these systems proved to be demanding, both with regard to atmosphere purity, to avoid oxygen pickup during sintering, and regarding removal of the oxygen introduced through the starting powders. By thermoanalytical studies it was shown that Cr prealloyed powders require significantly higher temperatures to remove the starting oxygen than Fe-C or classical alloy steel grades, as a consequence of the more stable oxides present.30
“Internal Getter” Effect in Powder Mixes
“Internal Getter” Effects in Prealloyed Systems
In prealloyed powders, typically produced by atomizing a suitably alloyed melt, the alloy elements are evenly distributed within the powder particles. All particles thus have virtually the same composition, at least with regard to the metallic constituents. Therefore, there should not be any heterogeneities of the oxygen affinity. In reality, however, also here oxygen transfer effects are observed.
If the sintering process is done in inert atmosphere, carbon is the main reducing agent, and it would be supposed that the first reduction peak, indicating removal of the surface oxygen, should emerge at 700–750°C, as encountered in Fe-C (see Fig. 3a). However, as evident from Fig. 6b, there is virtually no reduction (m28 = CO) peak at this temperature, but the first reduction stage occurs at about 1000°C, i.e. at a temperature level where reduction of Cr oxides typically occurs.
The MS graphs clearly show that annealing in Ar at 400°C does not change the reduction behaviour compared to the direct run in H2 done without anneal; in both cases there is a pronounced m18 (H2O) peak between 300°C and 500°C, with maximum closely below 400°C, indicating the reduction of the iron oxides present on the surface. The remaining oxides are reduced at much higher temperatures—above 900°C—with the maximum at about 1200°C. Once more carbon is the reducing agent. After annealing at 650°C, in contrast, the m18 peak in the 400°C range has almost completely disappeared, indicating that the iron oxide has already been transformed into more stable oxides—probably chromite (FeCr2O4)—even at this relatively low temperature, and without any soaking period.
This can be explained by assuming that in this temperature range, Cr is already sufficiently mobile in the iron lattice. It can diffuse, at least in a shallow zone, from the particle interior to the surface, where it reacts with the oxygen present. The driving force is the very negative Gibbs free energy of formation for the Cr oxide. Of course the diffusion distance of Cr is rather short—about 50 nm for 2 min at 640°C, according to extrapolated data from,38 but when considering that the iron oxide layer itself is only 6–7 nm thick, even the Cr from this shallow subsurface layer is sufficient to convert the iron oxide into a Cr based compound.
This can be regarded as the inverse process to that observed with Ti as mentioned above, which is in fact also an “internal gettering”. With Ti, the bulk metallic phase acts as the “getter” and oxygen is the diffusing element, migrating into the bulk. With the alloy steel, in contrast, the metallic constituent diffuses, migrating towards the surface. In both cases, however, the difference of the chemical potential between surface and bulk is the driving force for element transport.
Internal Getter Effects: Consequences for Industrial Practice
Internal getting is unwelcome from the practical side since it increases the temperatures required for oxygen removal and thus the sintering temperature. In industrial practice this means that more expensive furnaces may have to be used such as walking beam furnaces in place of mesh belt types. Therefore, avoiding this effect could be attractive also for the industry.
As has been shown, internal gettering requires exceeding a certain temperature threshold. For prealloyed powders, a temperature of 400°C is uncritical, and therefore using the reducing power of H2 in this temperature range to remove at least that part of the oxygen that is present as iron oxide is an option. Of course this oxygen fraction is typically 50% or less of the total oxygen content which implies that, even if this oxygen fraction is removed, still sintering at high temperatures is required to remove the remaining oxygen. Not surprisingly it has been shown that for Cr-Mo prealloyed steel grades, sintering at 1250°C and above is the most efficient way to obtain the excellent mechanical properties inherently possible with this alloy system. For the monotonic properties this has been described by Kremel et al.39 and for fatigue in Refs. 40 and 41. Hryha et al. have shown that when sintering at moderate temperatures, H2 in the atmosphere, as N2–H2 mix, is beneficial compared to N2 since part of the oxygen is removed early, but the differences between the sintering atmospheres tend to disappear at T > 1200°C.42
From the practical viewpoint it has to be considered that reducing the iron oxide at 400°C, in order to avoid the higher temperatures that result in internal gettering, is not so easy in industrial parts production. The temperature interval 400–600°C is exactly that in which lubricant burnout is done, and during this process the composition of the atmosphere is poorly defined, also oxygen containing fragments being present (the common lubricant EBS contains about 5 mass% oxygen). Karamchedu et al. have shown that in the laboratory, fairly complete delubrication can be attained at temperatures as low as 450°C44 which are uncritical temperatures with regard to internal gettering. However it remains questionable if such low temperatures are effective also for industrial sintering in which case several 100 kg of parts per hour have to be delubricated and sintered. In this case, higher delubrication temperatures seem to be inevitable.
In any case, for sintered parts from Cr- and Cr-Mo prealloyed powder compacts, high sintering temperatures are highly beneficial anyhow. The same holds for masteralloy-based mixes for which high temperatures are essential for attaining sufficient microstructural homogeneity.45 Therefore, although the internal getter effect may be unwelcome, its consequences are finally eliminated if high temperature sintering is applied.
For sintering of low alloy steels that contain admixed elements with high oxygen affinity—such as Cr, Mn and/or Si—the different oxygen affinity compared to the base iron has to be considered.
In powder mixes, the major proportion of the oxygen is typically introduced through the base iron powder. This oxygen is present as iron oxide on the particle surfaces and, after compaction, the oxide remains in the pressing contacts.
Superficial iron oxides are reduced at fairly low temperatures, but the gaseous reduction products tend to react with the alloy element particles, forming oxides that are much more difficult to reduce than the iron oxides, requiring typically temperatures >1000°C.
For prealloyed powders, a major proportion of the oxygen introduced into the compact through the starting powders is initially present as iron oxide, too. However, even at temperatures as low as 650°C and short times, they are transformed into more stable oxides that cannot be reduced by H2 at low temperatures, as can be iron oxide.
This means that internal gettering, and the resulting loss of reducibility of the oxides, occurs at much lower temperatures than assumed in the literature, in a temperature range typical for delubrication of pressed PM parts.
For the mixed systems, the critical temperature threshold for internal gettering is the higher one of two temperatures: either that of the first reduction stage forming gaseous compounds or the onset temperature of reactivity between the alloy element particles and these compounds
Masteralloy particles with passivating layers are less sensitive here since the second temperature threshold is higher.
For prealloyed powders the threshold is the temperature at which at least some diffusion of the alloy elements from the bulk to the powder surfaces is possible.
Internal gettering can to some degree be avoided by sintering in H2 or at least H2 containing atmospheres, the iron oxides being reduced at 300–500°C, but as stated above, in industrial practice this will interfere with the delubrication process.
In any case, both for Cr prealloyed systems and for mixed ones containing e.g. masteralloy powders, sintering at high temperatures −1250°C and above—is essential or at least beneficial for obtaining excellent mechanical properties. At these temperatures also the stable oxides formed by internal gettering are finally reduced.
Open access funding provided by TU Wien (TUW). This work was in part carried out within the international project “Höganäs Chair”, sponsored and organized by Höganäs AB, Sweden. The authors want to acknowledge financial support from the European Union through a Marie Sklodowska-Curie scholarship (Grant Agreement PIEF-GA-2013-625556) to one of the authors (R. de Oro Calderon).
- 1.S.-J.L. Kang, Sintering (Oxford, UK: Elsevier, 2005).Google Scholar
- 2.G.C.Kuczynski, ed., International Conference of “Sintering and Related Phenomena” (Notre Dame, IN: University of Notre Dame, 1965).Google Scholar
- 3.G.C. Kuczynski, Met. Trans. AIME 185, 169 (1949).Google Scholar
- 4.H.E. Exner, Grundlagen von Sintervorgängen (Berlin: Gebr Borntraeger, 1978) (in German).Google Scholar
- 5.J.E. Geguzin, Physik des Sinterns (Leipzig: VEB Deutscher Verlag für Grundstoffindustrie, 1973) (in German).Google Scholar
- 6.W. Schatt, Sintervorgänge (Düsseldorf: VDI-Verlag, 1992) (in German).Google Scholar
- 8.G. Petzow and W.J. Huppmann, Z. Metall. 67, 579 (1976) (in German).Google Scholar
- 10.S. Storchheim, Progr. Powder Metall. 18, 124 (1962).Google Scholar
- 11.H.-C. Neubing and G. Jangg, Met. Powder Rep. 42, 354 (1987).Google Scholar
- 12.G. Gille, G. Leitner, and B. Roebuck, Proceedings of the EuroPM’96 Stockholm (Shrewsbury UK: EPMA, 1996), pp. 195–210.Google Scholar
- 13.K.J.A. Brookes, Hardmetals and Other Hard Materials (East Barnet, UK: International Carbide Data, 1992).Google Scholar
- 14.M. Qian, Int. J. Powder Metall. 46, 29 (2010).Google Scholar
- 18.A.R. Glassner, The Thermochemical Properties of the Oxides, Chlorides, and Fluorides to 2500 K. U.S. Atomic Energy Commission Report ANL-5750, Washington, D.C. (1957).Google Scholar
- 19.I. Donaldson, M. Marucci, and B. Lindsley, Adv. Powder Metall. Partic. Mater. 2011, compiled by I.E. Anderson, T.W. Pelletiers, MPIF, Princeton, NJ Part 7 (2011), pp. 54–63.Google Scholar
- 20.J. Tengzelius, S.-E. Grek, and C.-A. Blände, Mod. Dev. Powder Metall. 13, 159 (1981).Google Scholar
- 21.A. Šalak, Powder Metall. Int. 18, 266 (1986).Google Scholar
- 22.A. Šalak, M. Selecka, and R. Bures, Powder Metall. Prog. 1, 41 (2001).Google Scholar
- 23.S. Banerjee, G. Schlieper, F. Thümmler, and G. Zapf, Mod. Dev. Powder Metall. 13, 143 (1981).Google Scholar
- 24.A.N. Klein, R. Oberacker, and F. Thümmler, Powder Metall. Int. 17, 71 (1985).Google Scholar
- 25.K. Ogura, S. Takajo, N. Yamato, Y. Maeda, and Y. Morioka, Met. Powder Rep. 42, 292 (1987).Google Scholar
- 26.S. Unami, K. Ogura, and S. Uemono: Proceedings of PM’98 World Congress Granada, vol. 3 (Shrewsbury, UK: EPMA, 1998), pp. 173–177.Google Scholar
- 27.I. Karasuno, K. Koshiro, M. Umino, and M. Ichidate, Horizons in Powder Metallurgy (Proceedings of PM’86 Düsseldorf), vol. 1 (Freiburg: Verlag Schmid, 1986) pp. 53–56.Google Scholar
- 28.B. Lindqvist, Proceedings of EuroPM2001 Nice, vol. 1 (Shrewsbury, UK: EPMA, 2001), pp. 13–21.Google Scholar
- 29.S. Berg and B. Maroli, Advances in Powder Metallurgy and Particulate Materials 2002 (Proceedings of PM2002 Powder Metall. World Congress, Orlando FL), Part 8 (Princeton, NJ: MPIF, 2002), pp. 1–14.Google Scholar
- 30.H. Danninger, C. Gierl, S. Kremel, G. Leitner, K. Jaenicke-Roessler, and Y. Yu, Powder Metall. 46(3), 1349, 2(3), 125 (2002).Google Scholar
- 32.C. Gierl-Mayer, H. Danninger, R. Oro Calderon, and E. Hryha, Int. J. Powder Metall. 51, 47 (2015).Google Scholar
- 33.C. Gierl-Mayer and H. Danninger, Powder Metall. Prog. 15, 3 (2015).Google Scholar
- 35.H. Karlsson, L. Nyborg, S. Berg, and Y. Yu, Proceedings of EuroPM2001 Nice, vol. 1 (Shrewsbury, UK: EPMA), pp. 22–27 (2001).Google Scholar
- 38.R.C. Weast, eds., Handbook of Chemistry and Physics, 67th ed. (Boca Raton, FL: CRC Press, Inc., 1989).Google Scholar
- 39.S. Kremel, H. Danninger, H. Altena, and Y. Yu, Powder Metall. Prog. 4, 119 (2004).Google Scholar
- 40.M. Dlapka, H. Danninger, C. Gierl, E. Klammer, B. Weiss, G. Khatibi, and A. Betzwar-Kotas, Int. J. Powder Metall. 48, 49 (2012).Google Scholar
- 42.D. Chasoglou, E. Hryha, and L. Nyborg, Proceedings of PM2010 World Congress & Exhibition, vol. 2 (Shrewsbury, UK: EPMA), pp. 3–12 (2010).Google Scholar
- 43.H. Danninger, A. Avakemian, C. Gierl-Mayer, M. Dlapka, and M. Grafinger, Proceedings of Euro PM2014, Paper-Nr. EP14038 (Shrewsbury, UK: EPMA, 2014).Google Scholar
- 44.S. Karamchedu, E. Hryha, and L. Nyborg, Powder Metall. Prog. 11, 90 (2011).Google Scholar
- 45.R. Oro, E. Bernardo, M. Campos, C. Gierl-Mayer, H. Danninger, and J.M. Torralba, Powder Metall. 59 (2015), in press.Google Scholar
Open AccessThis article is distributed under the terms of the Creative Commons Attribution 4.0 International License (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribution, and reproduction in any medium, provided you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons license, and indicate if changes were made.