powder metallurgy manufacturing of porous molybdenum ... · powder metallurgy progress, vol.13...

Powder Metallurgy Progress, Vol.13 (2013), No 3-4 93

POWDER METALLURGY MANUFACTURING OF POROUS MOLYBDENUM PREFORMS FOR Si INFILTRATION

E. Dadal, Ch. Gierl-Mayer, H. Danninger, L. Senčeková, K. Iždinský, F. Simančík

Abstract The present work focuses on preparation of porous Mo preforms that can be subsequently infiltrated with liquid Si to obtain an oxidation-resistant surface. Preforms pressed and sintered from commercial grade fine powders (typically < 10 µm) were found to be too reactive, because of the large specific surface, and the pore channels were rapidly choked by silicide formation, blocking further infiltration. Specimens from coarse powders with accordingly coarse open porosity were thus regarded more suitable. However, such powders are commercially available only as a rarity, at prohibitive cost. Therefore, coarse powders were specially prepared by presintering of compacts from the commercially available fine powders and afterwards crushing and sieving to the desired particle size. Powders ≤ 355 µm have thus been prepared, but their sintering activity left much to be desired. A suitable bimodal mix of coarse and fine powders, combined with appropriate sintering conditions, was found to result in well sintered preforms with clearly defined, infiltrable porosity. Keywords: high temperature materials, molybdenum, silicide, oxidation, infiltration

INTRODUCTION There are many materials that can be used at elevated working temperatures; these

are e.g. refractory metals or materials like graphite, carbides, oxides or nitrides; in many cases, however, the stability in oxidizing environment is insufficient. In many technical areas, progress is strongly restricted by the limits of current structural and protection materials. High-tech applications working in extreme conditions and environments such as aircraft and automobile engines, space vehicles, high temperature or high pressure fusion or chemical reactors require maximum performance and thus revolutionarily improved materials able to withstand complex loading conditions. For such applications, typically many loading factors are combined, such as e.g. high temperature, corrosive (oxidative) attack and multiaxial dynamically changing forces. Furthermore, requirements like availability of techniques for cost efficient production of net shape (also very complex) structural parts should be fulfilled [1,2]. Graphite and other carbon based materials are widely used as high temperature structural materials due to their advantages of dimensional stability, easy machinability, low weight and relatively low cost. However, they degrade rapidly at temperatures above 800°C in oxidizing atmospheres. The major limitation to the use of diborides and carbides of the early transition metals (e.g., Zr, Hf, Ta), which are candidates for the use in extreme environments because of their thermal and chemical Ewa Dadal*, Christian Gierl-Mayer, Herbert Danninger, Vienna University of Technology, Faculty of Technical Chemistry, Institute of Chemical Technologies and Analytics, A-1060 Vienna, Austria Lucia Senčeková, Karol Iždinský, František Simančík Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Račianska 75, SK-831 02 Bratislava, Slovak Republic *now: Traktionssysteme Austria GmbH, A-2351 Wiener Neudorf, Austria


stability, is a lack of an affordable, reproducible manufacturing process. Oxides are stable in an oxidizing environment but they suffer from lack of thermal shock resistance, low fracture toughness, and poor creep resistance at elevated temperature. Ni based superalloys still remain as the most reliable and widely used material systems for many high temperature applications. They exhibit excellent creep strength, oxidation resistance, and fracture toughness [3]. However, their fundamental limitation is their melting point. Since advanced superalloys melt at temperatures in the range of 1350°C, significant strengthening can be obtained only at temperatures below 1150°C, which is as much as 150 K lower than requirements defined by designers of efficient and clean aircraft or modern industrial turbines [4,5]. The next group are the refractory metals which have high thermal shock resistance and high fracture toughness. Mo and Nb metals, for example, satisfy many of the requirements for engine applications. One crucial disadvantage is however that both metals suffer from poor oxidation resistance. The oxide layer that forms on Nb, - Nb2O5 - does not protect the metal from further oxidation. For Mo it is the MoO3 oxide which is formed, but it is volatile above about 700°C [6]. They are also limited by creep resistance at elevated temperatures. In addition, most of them have high densities and cost.

For applications at higher temperatures in air or other oxidizing atmospheres it is necessary to protect molybdenum against oxidation. One possibility to protect the material is to use a protecting layer e.g. of MoSi2 on the Mo bulk material. MoSi2 possesses excellent oxidation resistance in oxygen atmosphere between 1200°C and 1700°C due to formation of a thin protective SiO2 layer at the surface of the MoSi2 bulk. This protective ability is due to the volatility of MoO3 which evaporates at 750°C, leaving behind an amorphous SiO2 film which is embedded in the MoSi2 matrix and will protect the underlying metals from oxidation from 600°C to above 1700°C (5MoSi2 + 7O2 --> Mo5Si3 +7SiO2). [7,8]. The protection of MoSi2 layers depends on the oxygen partial pressure of the gas atmosphere and temperature. At temperatures below 1200°C or at reduced pressure the protective properties of the MoSi2 layer suffer [9]. A considerable problem is the “pesting” of MoSi2

which is the oxidation of MoSi2 at lower temperatures [10]. At temperatures between 400 and 600°C pesting can result even in structural disintegration of the material [10]. However, below 600°C decomposition may occur due to oxidation when the formation of MoO3 competes with that of SiO2 (2MoSi2 + 7O2 2MoO3 + 4SiO2) [7, 8]. The oxidation resistance of molybdenum can also be increased by infiltration of porous MoSi2 with an SnAl-alloy [11]. The combination of brittle silicides with a low melting alloy increases the thermal shock resistance and the self healing capacity of propagating cracks. However, the low melting point of the SnAl alloy disqualifies it as high temperature oxidation resistant material. Some other references report a better oxidation resistance of Mo-Si-B alloys [12,13] where the (Mo rich) Mo-Si-B alloy forms an adherent borosilicate layer during high temperature oxidation. By adjusting the Si content, the high temperature oxidation resistance can be balanced by the formation of a passivating SiO2 rich layer. A too high amount of Si however will impair the oxidation resistance at low temperatures. In a further work a hypothesis was presented for the mechanisms of better oxidation resistance by the addition of rare earth oxides which reportedly results in lowering the oxygen diffusion at the grain boundaries. It could be shown that the oxidation could be slowed down but was not avoided completely [9].

For high temperature applications, MoSi2 coatings offer a good protection against oxidation by forming stable SiO2 coatings [8,11,13,15-19]. It should however be mentioned that such MoSi2 layers offer problems with adhesion on the bulk material by cracking. During oxidation experiments such cracking and spalling of the MoSi2 layer was observed, too [16]. One possible way to produce such Mo specimens with well adhering MoSi2


surface is by infiltration of at least partially porous molybdenum bodies with Si, to generate an interlocking interface. For sufficient infiltration it is necessary to provide larger open pores for silicon to penetrate and cover the whole Mo body while retaining a graded transition from Mo to the silicide. However, for preparation of sintered bodies with sufficiently large pores the particle size of Mo is the most important factor. With increasing average size of the Mo powder used, the size of the open pores increases too. Due to the fact that coarse Mo powder is barely available on the market, one possible way might be to reduce Mo oxide in a humid H2 atmosphere in a high powder bed [19-22]. Experiments with W, which is chemically similar to Mo, have demonstrated that with increasing powder bed height and reduction temperature, the water vapour content in the system increases, too [23]. With increased water vapor content and increased temperature, more volatile WO2(OH)2 is formed which is adsorbed on energetically favourable crystal faces, forming larger grains as a result. It was also reported [9,24-26] that the addition of impurities such as LiOH to W oxides promotes grain growth during reduction while evaporating without Li residues. In the first reduction steps to WO2 no morphological differences between doped and undoped tungsten could be seen. In the last reduction step from WO2 to W an increase in the reaction rate could be observed. That results in a higher reduction rate by the alkaline metals, which increases on the other side the water production. By the increase of water also the number of nuclei is decreasing. This results in a smaller amount of new nuclei but of a larger size of the crystals [24,27].

EXPERIMENTAL Various experiments were carried out to prepare coarse Mo powders. In the first

two approaches it was tried to prepare coarse powder by the reduction of MoO2 powder under different conditions. By the variation of the powder bed height or the addition of LiOH it was tried to increase the resulting powder size. In the third experiment some presintered specimens pressed from commercial fine Mo powder were crushed, and the resulting powder was annealed.

In the first reduction approach, MoO2 powder (supplied by Chemie Metall, Bitterfeld, Germany) was reduced at three different temperatures, at 900°C, 1000°C and 1100°C, in flowing H2 (flow rate 2 Nl/h; 5.0 purity) for 2 h isothermally. For all reduction experiments an electrically-heated pushtype tube furnace with heat resistant steel muffle with 70 mm inner diameter and Kanthal heating elements was used. In a further experiment MoO2 was reduced at the same conditions but in a higher powder bed. The unreduced MoO2 was characterised by SEM and the images are shown in Figs.1a and 1b. The particles are angular and platelet-shaped and are in the size range of about 4 µm.

Fig.1. SEM images of MoO2 powder (Chemie Metall) at different magnifications.


In the first tests the MoO2 powder was reduced in a dry H2 (5.0 purity) atmosphere. In a second experiment MoO2 was reduced in humid H2 atmosphere. To establish a humid gas atmosphere in the furnace tube during all reduction experiments, the carrier gas H2 was directed through a gas-washing bottle filled with distilled water. The H2 gas had to pass the water, absorbed water vapour and was directed into the furnace tube. The height of the powder bed was chosen as 20 mm. The oxygen content of the powders was measured by hot fusion analysis (LECO TC400), and their morphology was characterised in the scanning electron microscope (FEI Quanta 200).

In a further experiment MoO2 was mixed with 250 ppm Li (as LiOH; from Merck) before reduction. For better homogenisation the LiOH was dissolved in distilled water. The MoO2 was poured into the solution and stirred for about 10 min mechanically. The suspension was then dried in a drying box at 120°C for 24 h until the powder formed a dry and stable cake which was crushed and sieved through a 355 m sieve. Afterwards the Li doped oxide powder was reduced in a tube furnace at 1100°C for 2 h in a humid H2 atmosphere with a 40 mm powder bed height.

In the third experiment a very coarse Mo powder was prepared by pressing commercial grade Mo powder (standard “≤63 µm” Mo from Chemie Metall, in reality <10 µm) with 200 MPa to rectangular bars (Charpy bars ISO 5754). The green compacts were first deoxidized and presintered in a pushtype furnace at 1100°C for 2 h in H2 (2 Nl/h gas flow; 5.0 purity) and then sintered at 1500°C for 2 h in vacuum. Afterwards the specimens were crushed in liquid N2 and milled in a knife mill for 10 s. The crushed powder was sieved through a 355 µm mesh sieve. The coarser residue was milled again until all of the powder had attained the desired particle size of ≤355 µm. The powder was then characterised by SEM.

Finally, a bimodal powder mix was prepared by mixing a commercial coarse powder grade (Goodfellow, <355 µm) with a fine grained one (once more from Chemie Metall; ≤63 µm Mo, in fact <10 µm) in the ratio 50:50 % for 5 h in a tumbling mixer. From this powder mixture Charpy bars with 55*10*10 mm were pressed at 200 MPa. The specimens were presintered at 1100°C for 2 h in H2 (2 l/h; 5.0 purity) in the pushtype furnace followed by sintering at 1800°C in vacuum for 2 h. For this experiment, a KCE 100 kN vacuum hot press was used for pressureless sintering, i.e. without axial pressure. Subsequently, infiltration tests with liquid Si were performed, and the resulting bodies were characterized by metallography and microanalysis.

RESULTS

Reduction experiments MoO2 was reduced in dry H2 in the first experiment and in humid H2 in a second

experiment and at three different temperatures. In one further run MoO2 was reduced in humid H2 atmosphere at 1100°C in a higher powder bed. For MoO2 powder reduced in dry H2 atmosphere the measurement of the oxygen content showed that with increasing temperature up to 1000 °C the oxygen content is decreasing to 0.1 ± 0.1 mass% and with further increasing temperature up to 1100 °C the oxygen content is increasing to 0.2 ± 0.1 mass%. The results for the reduction experiments humid H2 atmosphere showed that with increasing reduction temperature the oxygen content decreases. After reduction at 900°C the oxygen content was about 1.0 ± 0.1 mass%, after 1000°C it decreased to 0.7 ± 0.1 mass%. After reduction at 1100°C the content of oxygen further decreased to 0.5 ± 0.1 mass% and with higher powder bed it could be decreased even to 0.4 ± 0.1 mass%.


In Figure 2 the Mo powder after reduction in dry H2 atmosphere is shown. It can be seen very well that the shape of the resulting particles is changing with the reduction temperature. After reduction at 900°C the shape of the Mo particles remained plate-like. By increasing the temperature the edges of the particles were progressively rounded, which can be observed very well in Fig.2b - after 1000°C reduction. After reduction at 1100°C round particles of different sizes can be recognised. The largest particles are about 5 µm in diameter.

Fig.2a. 900°C Fig.2b. 1000°C

Fig.2c. 1100°C

Fig.2. Mo powder obtained by reduction of MoO2 in dry H2 atmosphere at different temperatures.

In Figure 3 the shape and size of the powder particles after reduction in humid H2 atmosphere can be observed very well. Comparing the unreduced MoO2 powder with the powders reduced at different temperatures it is evident that also in this case the particles lost their plate-like shape with increasing reduction temperature. After reduction at 900°C the particles were still angular and more like platelets. At 1000°C the plates are converted more to octahedrons. The powder reduced at 1100°C consists out of fairly round particles with varying size. Some particles were very small, about 1 µm in diameter, and some others had grown to ~ 4 µm. After reduction at 1100°C but in a higher powder bed all particles were round and of similar size of about 4 µm; however, the expected coarsening hat not occurred.

In a further experiment MoO2 was mixed with 250 ppm Li (as LiOH) in distilled water, dried and reduced afterwards at 1100°C for 2 h in H2. It can be observed in Fig.4 that the powder does not coarsen after reduction, as was originally expected based on the expected


analogy with W. The shape and size of the particles is similar to that of the powder without LiOH doping and reduced at 1100°C for 2 h in the same atmosphere (Fig.3b). This indicates that the effect of alkaline elements – at least of Li – on the particle size during reduction of Mo oxides is much less pronounced that in the case of tungsten oxides.

Fig.3a. 900°C Fig.3b. 1000°C

Fig.3c. 1100°C Fig.3d. 1100°C, higher powder bed

Fig.3. Mo powder obtained by reduction of MoO2 in humid atmosphere at different temperatures.

Fig.4a. doped MoO2 before reduction Fig.4b. Mo reduced from Li doped oxide

Fig.4. Reduction experiment with 250 ppm Li doping.


Bimodal powders In the final approach some very coarse Mo powder was prepared from the fine Mo

grade from Chemie Metall as described above, i.e. by pressing, sintering, crushing and milling. Different particle sizes of the resulting powder are depicted in Fig.5a. Many particles are up to 355 µm in size but also a lot of very small particles remain, indicating disintegration of some of the coarser particles during comminution. At higher magnification (Fig.5b) many sintering bridges can be observed within one powder particle, which was expected due to the very high sintering temperatures.

Fig.5a. overview Fig.5b. detail from coarse particle

Fig.5. SEM images of coarse Mo powder - ≤ 355 µm.

Pressed and sintered specimens prepared from the coarse powder showed the desired coarse porosity and low specific surface, but the strength of the sintering contacts – and thus of the sintered bodies themselves - was very poor, as a consequence of the low specific surface and resulting low sintering activity of the coarse powders.

Fig.6a. overview Fig.6b. detail

Fig.6. Section of Mo sample prepared from bimodal powder (commercial coarse and fine grained Mo powders 50:50), presintered 2 h at 1100 °C in H2; sintered 2 h at 1800 °C in

vacuum .

To eliminate this deficiency, a bimodal powder mix was prepared by mixing a commercial grade coarse powder (which had a more clearly defined particle size distribution than the self-prepared grade), with fine grained standard Mo in the ratio of


50:50 %. In Figure 6 the cross section of such a pressed and sintered sample is shown. A lot of interconnected pores can be seen. The fine grained powder acts as “binder” for the coarse particles, thus compensating for the low sintering activity of the plain coarse powders. This results in the desired microstructure with fairly low specific surface - and thus controllable reactivity with liquid Si - and on the other hand sufficient strength of the necks between the coarse particles to prevent disintegration during infiltration.

After infiltration with liquid Si, most pores were closed by the formation of molybdenum silicides. Disintegration of the sintered Mo bodies, which proved to be a major problem with porous Mo specimens from fine powders, could in fact be avoided due to the strong interparticle bonding. In Figure 7 the cross section of the infiltrated sample can be observed. Only a few pores remained after infiltration. In Figure 7b, the persistent sintering bridges between the particles can be observed.

Fig.7a. Transition area infiltrated zone-core Fig.7b. detail from infiltrated zone

Fig.7. Section of Mo specimen as in Fig.6 after infiltration with liquid Si.

Fig.8. Specimen as in Fig.7; SEM image and EDAX analysis of Si infiltrated area .

By EDAX analysis it was possible to describe the different areas of the sample that are shown in Fig.8 in different colours. The darkest areas with ~36 wt.% Si are MoSi2, the lighter ones with ~16 wt.% Si are Mo5Si3, and the lightest particles, which form the core of some large grains, show a very low Si content, with 0.3 wt.% Si (which indicates more or less plain Mo). It can be observed that the core of each particle is completely surrounded by Mo5Si3 and MoSi2 as outer shell.


CONCLUSIONS For protecting Mo in oxidizing environment at elevated temperatures, silicide

coatings were applied by infiltrating a porous Mo body. It could be learned from the experiments that the infiltration of porous Mo bodies with Si is a challenge. The first tests showed that the fine pore structures obtained when using commercial fine Mo powder grades are not suitable, resulting in fast blocking of the infiltration process by silicide formation and in excessive exothermic reactions, with resulting disintegration of the preform. Preforms with larger open pores were expected to enable a faster infiltration before reaction (of Si with Mo) and closing of the pores. Therefore, coarser Mo powders were required, which should result in lower reactive surface, but these are not commercially available, at least not at acceptable cost. Several routes were tried to obtain such coarse Mo powders, using experience derived from W powders.

Since it is well known that reduction of W oxides in dry and humid H2 results in coarser powder, MoO2 was reduced at 900°C / 1000°C / 1100°C (and in higher powder bed) in dry and humid H2. In dry H2 atmosphere the lowest oxygen content was achieved by sintering at 1000°C. Above this temperature the water vapour pressure changed and the reduction of MoO2 was not as favoured as before. In the SEM images of Mo powder reduced in humid H2 atmosphere could be seen that even at the highest reduction temperature the particles did not exceed 5 µm in size. The oxygen content of the Mo powder obtained decreased with higher reduction temperature and, surprisingly, also with higher powder bed. With the reduction temperatures also the shape of the particles changed; at lower temperatures the particles were more like flat plates, resembling the shape of the original oxide particles, while the fully reduced powder (after reduction at 1100°C in a higher powder bed) consisted of round particles. However, the size of the round particles was only about 4 µm, which means that reduction in humid atmosphere was only moderately effective.

Also by doping the MoO2 with 250 ppm LiOH and reduction at 1100°C in H2 the particle size could not be increased, which is in marked contrast to the results obtained with W as described in the literature. The shape and size (~ 4 µm) of the particles are similar to MoO2 reduced at 1100°C in a higher powder bed without LiOH doping, i.e. the effect of humidity is at least as pronounced as that of Li doping, but both are significantly less pronounced than with W.

Finally very coarse Mo powder was prepared by crushing of sintered Mo specimens (sintered at 1500°C for 2 h; vacuum) prepared from fine grained Mo. This coarse powder was used for preparation of porous preforms. Specimens prepared exclusively from coarse powder were found to be very fragile, due to the very low sintering activity of this powder. However, by mixing coarse with fine grained powder at a ratio of 50:50 % and subsequent pressing and sintering, specimens with acceptable strength were obtained. A network of percolating coarse pores can be observed while the density of the material remains still high at 7.9 g/cm3 but with 22.5 % open porosity. Infiltration tests showed that this bimodal porosity in fact is infiltrable by liquid Si while disintegration of the preform is suppressed.

Acknowledgement The present work has been performed within the project SILTRANS, funded by

the EU through the 7th European framework programme. The authors want to thank the project partners for their cooperation.


REFERENCES [1] Bewlay, BP., Jackson, MR., Zhao, JC., Subramanian, PR.: Metall. and Mater. Trans. A,

vol. 34, 2003, no. 10, p. 2043 [2] Schneibel, JH.: Beyond Nickel-Base Superalloys. Oak Ridge : Metals and Ceramics

Division Oak Ridge National Laboratory, 2008 [3] Raisoon, G.: Powder Metall., vol. 51, 2008, no. 1, p. 10 [4] Baldan, A.: Journal of Material Science, vol. 37, 2002, p. 2379 [5] Nabarro, FRN.: Metall. Trans. A, vol. 27, 1996, p. 513 [6] Perepezko, J., Sakidja, R.: Adv. Eng. Mater., vol. 11, 2009, p. 892 [7] Byun, JY., Yoon, JK., Kim, GH., Kim, JS.; Choi, CS.: Scripta Materialia, vol. 46,

2002, p. 537 [8] Lassner, E., Püschel, R., Benesovsky, F.: Planseeberichte Pulvermet., vol. 16, 1968, p.

91 [9] Chao, PJ., Priest, DK.: J. Less-Common Metals, vol. 2, 1960, p. 426

[10] Chou, TC., Nieh, TG.: JOM, vol. 45, 1993, p. 15 [11] Stecher, P., Lux, B.: J. Less-Common Metals, vol. 14, 1968, p. 407 [12] Zamoum, F., Cartigny, Y., David, N., Fiorani, JM., Podor, R., Vilasi, M. In: Proc. 16th

Plansee Seminar. Vol. 1. Reutte, 2005, p. 883 [13] Sharif, AA.: J. Mater Sci, vol. 45, 2010, p. 865 [14] Jehanno, P., Heilmaier, M.: Powder Metall., vol. 51, 2008, no. 2, p. 99 [15] Melsheimer, S., Rahmel, A., Schütze, M. In: Proc. 13th Plansee Seminar. Vol. 1.

Reutte, 1993, p. 939 [16] Cox, AR., Brown, R.: J. Less-Common Metals, vol. 6, 1964, p. 51 [17] Yao, Z., Stiglich, J., Sudarshan, TS.: J. Mater. Eng. Perform, vol.. 8, 1999, p. 291 [18] Petrovic, JJ.: Materials Sci. Engrg, vol. 192-193, 1995, p. 31 [19] Kim, SH., Kim, DG., Seo, YI., Kim, YD.: Rev. Adv. Mater. Sci., vol. 28, 2011, p. 141 [20] Haber, J.: J. Less-Common Metals, vol. 57, 1977, p. 243 [21] Tuominen, SM., Carpenter, KH.: J. Metals, vol. 32, 1980, p. 23 [22] Schulmeyer, W., Ortner, HM.: Int. J. Refractory Metals and Hard Materials, vol. 20,

2002, p. 261 [23] Argent, BB., Milne, GJC.: J. Less-Common Metals, vol. 2, 1960, p. 154 [24] Haubner, R.: Untersuchung der Einflüsse von geringen Konzentrationen an

Alkalimetallen, Aluminium, Eisen, Nickel, Phosphor und Silizium auf die Reduktion von WO3 zu Wolfram mit Wasserstoff. PhD thesis. Wien : Technische Universität, 1984

[25] Danninger, H., Jangg, G., Lassner, E., Lux, B.: High temp. – high press., vol. 13, 1981, p. 541

[26] Danninger, H., Pisan, W., Jangg, G., Lux, B.: Z. Metallkunde, vol. 74, 1983, p. 151 [27] Haubner, R., Schubert, WD., Lassner, E., Lux, B. In: Proc. 11th Int. Plansee Seminar.

Vol. 2. Reutte, 1985, p. 69

powder metallurgy manufacturing of porous molybdenum ... · powder metallurgy progress, vol.13...

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