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Ductilization of Mo–Si solid solutions manufactured by powder metallurgy
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Acta Materialia 57 (2009) 3895–3901 www.elsevier.com/locate/actamat
Ductilization of Mo–Si solid solutions manufactured by powder metallurgy H. Saage a,*, M. Kru¨ger b, D. Sturm b, M. Heilmaier c, J.H. Schneibel b, E. George d, L. Heatherly d, Ch. Somsen e, G. Eggeler e, Y. Yang f b
a Faculty of Mechanical Engineering, University of Applied Sciences, Landshut, D-84036 Landshut, Germany Institut fu¨r Werkstoff- und Fu¨getechnik, Otto-von-Guericke-Universita¨t Magdeburg, D-39016 Magdeburg, Germany c Fachbereich Materialwissenschaft, TU Darmstadt, D-64287 Darmstadt, Germany d Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA e Institut fu¨r Werkstoffe, Ruhr-Universita¨t Bochum, D-44801 Bochum, Germany f Computherm LLC, 437 S. Yellow Stone Drive, Madison, WI 53719, USA
Received 18 March 2009; received in revised form 20 April 2009; accepted 23 April 2009 Available online 22 May 2009
Abstract Mo–1.5 at.% Si alloys with additions of either Y2O3 or Zr were manufactured by mechanical alloying. The Y2O3 particles reduced the grain size and increased the room temperature strength, but did not alleviate the brittleness of previously investigated Mo–1.5 at.% Si without Y2O3. Additions of Zr, on the other hand, resulted not only in a fine grain size and an extremely high bend strength (2 GPa), but also in limited bend ductility at room temperature. Zr additions are seen to be beneficial for three reasons. First, Zr reduces the grain size. Second, Zr getters detrimental oxygen by forming ZrO2 particles (which in turn help to pin the grain boundaries). Third, in situ Auger analysis shows that Zr reduces the concentration of Si segregated at the grain boundaries. This is thought to enhance the grain boundary cohesive strength and thus leads to the observed ductility. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mechanical alloying; Bending tests; Molybdenum silicides; Ductility; Microstructure
1. Introduction A family of three-phase Mo–Si–B alloys pioneered by Berczik [1] is considered as a potential new high-temperature material class for applications in air, since it combines the excellent high-temperature strength of Mo-base alloys (such as TZM: Mo–0.5Ti–0.08Zr–0.03C, wt.%) and the beneficial oxidation resistance of silicide phases. The two intermetallic phases Mo3Si and Mo5SiB2 (T2), which form in addition to the Mo(Si) solid solution (‘‘a-Mo”) matrix phase, improve the high-temperature oxidation resistance since they provide the Si and B necessary for forming a coating of a dense, protective SiO2(B) glass [2].
*
Corresponding author. Tel.: +49 871 506 219; fax: +49 871 506 9 219. E-mail address: [email protected] (H. Saage).
Mo possesses a body-centred cubic (bcc) crystal structure and is characterized by a brittle-to-ductile transition temperature (BDTT; see e.g. [3]) which is accompanied by a change from a mostly intergranular brittle fracture mode to mostly transgranular above the BDTT. The transition temperature itself depends on the purity of Mo, which tends to fail in a brittle manner by intergranular fracture at room temperature in the presence of even small amounts of interstitial impurities (50 wt.ppm) such as oxygen [3,4]. This is attributed to the weakening of the grain boundaries caused by this element. The solubility of Si in Mo depends on temperature but can be as high as 2 at.% at 1400 °C, the designated maximum application temperature of Mo–Si–B alloys. Therefore, in a previous work [5] we studied the mechanical behaviour of high-purity Mo(Si) solid solutions with Si contents ranging between 0.1 and 1 wt.% (i.e. about 0.3–3.3 at.%) as a reference for
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.04.040
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the matrix phase of three-phase Mo–Si–B alloys. As an important outcome we could demonstrate that while Si is one of the most effective solid solution strengtheners for Mo [3], it also strongly promotes the tendency towards brittle intergranular failure [5]. It is well known that the ductility of molybdenum can be improved by reducing its grain size [3]. Thus, one of the approaches to ductilize Mo(Si) solid solutions in the present work is to add Y2O3, a thermodynamically very stable oxide, which reduces grain boundary mobility, and therefore stabilizes a small grain size during powder metallurgical (PM) manufacturing. However, this may have a detrimental impact on the creep performance at higher temperatures at which the material is intended to be used. Another approach is to strengthen the grain boundaries by reducing the oxygen concentration at the grain boundaries by adding carbon [4]. Therefore, a second and possibly more viable way is to find suitable microalloying additions which improve the ductility by reducing the tendency of solute atoms to segregate to the grain boundaries and embrittle them. Among the possible candidates which have been identified through a theoretical approach [6], Zr has been shown recently to improve the fracture toughness of Mo–12Si– 8.5B (at.%) [7]. According to Miller and Bryhan [8], who measured the distribution of elements at and close to the grain boundaries by atom probe tomography, the basic mechanism for the improvement of the grain boundary fracture stress of a welded Mo-base alloy containing small amounts of Zr, C and B was found to be the substitution of O and possibly N by these additions at the grain boundaries. Therefore, in our second approach we use ZrH2, which decomposes into its elements during manufacturing at high temperatures, in order to ductilize Mo(Si) solid solutions prepared by PM. 2. Materials and methods The maximum solubility of Si in Mo at the targeted application temperature of at least 1200 °C is about 1.5 at.% (estimated from extrapolation of the respective ternary phase diagram) [9]. Thus, elemental powders of Mo and Si of 99.95% and 99.9% purity, respectively, were mixed to obtain this target composition (compositions are given in at.% unless stated otherwise). About 0.7 vol.% of Y2O3 powder or about 1 wt.% ZrH2 powder, which is reduced to about 1 at.% Zr upon sintering, were added to obtain two additional compositions. Mechanical alloying (MA, a high-energy milling process [10]) was carried out for 20 h under a protective (argon) atmosphere in a planetary ball mill with a rotational speed of 200 rpm and a powder to ball weight ratio of 1:13. In the following all alloy designations without a prefix refer to alloys where MA was used to prepare the powders. In some cases PM is used as prefix instead which refers to alloys prepared from blended powders (not MAed). The detailed processing of these alloys is described elsewhere [5]. After milling
the powders were cold isostatically pressed at 200 MPa and sintered in hydrogen at 1500 °C to reduce the impurity content (mainly O, N, H). Depending on the preparation technique the powders respond differently to this treatment. Simple mixtures of the elements Mo and Si have impurity levels of below 5 wt.ppm [5], whereas the MAed powders studied here possess 200–500 wt.ppm oxygen after sintering, which remains at this level upon further consolidation. As a final step the sintered specimens were further consolidated by hot isostatic pressing (HIP) at 1500 °C and 200 MPa for 4 h. Electrodischarge machining was used to prepare bending specimens with a rectangular cross-section of about 3.8 3 mm and a length of 45 mm from the HIP billets and final surface quality was attained by grinding. The mechanical properties were assessed by three-point bending tests with a span of 40 mm in a temperature range of room temperature–1093 °C using a Zwick electromechanical testing machine equipped with a Maytec furnace at a crosshead speed of 0.01 mm min 1. At temperatures above 538 °C a gas mixture of argon with 2% hydrogen was used for the bend tests to prevent oxidation. Bend deformation was continuously monitored by Al2O3 rods attached to the centre of the lower (tensile) side of the specimens and to the lower pushrod, respectively. The stresses and (elastic) strains on the tensile side of the bend specimen were calculated utilizing elastic beam theory [11], which is sufficient to detect the BDTT, i.e. the temperature above which plasticity occurs. With this configuration even small deviations from the elastic behaviour could clearly be resolved. Specimens with the composition Mo–1.5Si and Mo– 1.5Si–1Zr were broken in situ in the ultra-high vacuum of an Auger microscope (PHI 680 SAM) and their fracture surfaces were analysed for possible segregation of the alloying elements at the grain boundaries. From secondary electron (SE) scanning electron microscopy images of the fracture surfaces, median values of the grain size were determined by directly measuring at least 250 grains. 3. Results 3.1. Mechanical alloying and consolidation Mechanical alloying of the Mo–1.5Si powder mixtures is accompanied by a significant microstructural refinement. ˚ durThe Mo domain size is reduced to values below 20 A ing milling and, typical for brittle–ductile-type powder mixtures [10,12], the additional phases Y2O3 and ZrH2 are first crushed into smaller pieces and subsequently introduced into the more ductile Mo powder particles to form fine, homogeneously distributed dispersions. Due to the substitution of Mo by smaller Si atoms the lattice parameter is reduced during MA and the dissolution of 1.5% Si is complete after milling for 20 h [5,13]. The ZrH2 decomposes during consolidation by releasing the hydrogen. Some of the Zr is thought to remain in solid solution, while the
H. Saage et al. / Acta Materialia 57 (2009) 3895–3901
remainder reacts with residual oxygen to form ZrO2 particles in the grain interior and at the grain boundaries (Fig. 1). The Si remains in solution after consolidation (HIPing) at 1500 °C. As shown in Table 1 the grain size after consolidation depends on both the composition as well as the processing conditions. While the median value of the grain size of the Mo(Si) solid solution obtained from MAed powder is about 3.2 lm, the addition of 1% Zr reduces the grain size to 1.7 lm. Alloying with Y2O3 dispersoids led to the lowest grain size, less than 0.5 lm, after consolidation. For comparison, specimens consolidated from simply mixed Mo– Si powders of our previous study [5] yielded mean grain sizes which were larger by more than one order of magnitude.
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tional Y2O3 dispersoids are brittle at temperatures up to 538 °C. However, both alloys show some plastic deformability at 816 °C. In contrast, the alloy with ZrH2 addition already reveals limited plastic strain at room temperature at a very high strength level of around 2 GPa. Moreover, this alloy reaches a plastic strain of about 1.4% at 538 °C while exceeding the strength level of 2 GPa. As expected, at even higher temperatures the strength of this alloy decreases while the plastic deformability increases. Fracture of the Mo(Si) solid solution occurs predominantly along the grain boundaries with a fraction of intergranular failure of about 93%, with the remainder showing transgranular fracture. In contrast, the Zr-containing alloy shows less intergranular fracture (80%, see Fig. 3), whereas fracture of the Y2O3-containing alloy occurs almost 100% along the grain boundaries at room temperature. For comparison, on the basis of tensile tests, we have shown [5] that the addition of only about 0.3 % Si leads already to brittle intergranular fracture at room temperature, while nominally pure Mo almost reaches 15% plastic strain before failure and reveals a transcrystalline fracture mode (see Table 1 and Fig. 4).
3.2. Temperature dependence of strength and ductility Stress–strain (calculated for the outer tensile fibre) curves for the alloys tested at room temperature and at 538 °C are compared in Fig. 2. Values for the yield strength, maximum bend strength and plastic strain before failure for all test temperatures are given in Table 1. The MAed Mo–1.5 Si solid solution and the alloy with addi-
3.3. Auger analysis on the fracture surface
Fig. 1. Bright-field TEM image (two beam condition) showing fine dispersed ZrO2 particles in the grain interior and the pinning of dislocations.
In Fig. 5a Auger spectra of the grain boundary regions (intergranular fracture areas) and the grain interior (transgranular) are compared for PM Mo–0.3Si as well as for MA Mo–1.5Si and MA Mo–1.5Si–1Zr solid solution. Unfortunately, it was not possible to reliably compare the oxygen concentration at the grain boundaries. Also, because of peak overlap with Mo, the Zr concentration could not be determined. The only obvious difference of the spectra taken from the grain boundaries and the grain interior is in the size of the Si peak at 91 eV. While this peak is pronounced for the grain boundaries, the transgranular fracture area (grain interior) shows, consistent with the low bulk Si concentration, only a small peak. Quantitative analyses of the Si concentration on the basis of several Auger spectra shown in Fig. 5b confirm that the Si concentration is reduced at the grain boundaries of the Zr containing Mo–1.5Si solid solution, as compared to Mo(Si) alloys without Zr. Hence, Si is strongly segregated at the grain boundaries in all binary Mo–Si alloys
Table 1 Different parameters of the MAed materials investigated here. In addition, tensile test values for two PM Mo–Si solid solutions which were taken from Ref. [5] are also shown. Alloy
Median of the Yield strength in MPa grain size in lm
Maximum bend or tensile strength Plastic strain in % in MPa
RT 538 °C 816 °C 1093 °C RT
538 °C
816 °C
1093 °C
RT 538 °C 816 °C 1093 °C
Mo–1.5Si (MAed) 3.2 Mo–1.5Si + 0.7 vol.% Y2O3 (MAed) 0.45 1.7 Mo–1.5Si + 1 wt.% ZrH2 (MAed)
– – –
– – 1950
920 960 1040
– 67 (3%) 385
– – – – 0.1 1.4
0.3 1.2 0.9
– >3% >3%
Mo–0.3Si (PM) Mo–1.5Si (PM)
– –
166 –
150 331
219 296
– –
27 3.3
90 0.6
54 35
– 48 305
516 1581 2001
731 1804 2075
953 991 1203
132 275
– –
377 –
365 541
32 –
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Fig. 2. Stress–strain curves (outer fibre) of pure Mo–1.5Si and Mo–1.5Si with Y2O3 and ZrH2 addition at (a) room temperature and at (b) 538 °C, (c) 816 °C and (d) 1093 °C.
but to a much lesser extent in the Zr-containing alloy, indicating that the presence of Zr leads to depletion of Si at the grain boundaries. 3.4. Calphad-type thermodynamic modelling of the Mo–Si– Zr ternary system The question arises whether the reduction of the Si grain boundary concentration due to Zr could be caused by a reduction in the bulk solubility in the presence of Zr. Therefore, the effect of Zr addition on the solubility of Si in Mo was studied by Calphad-type thermodynamic modelling of the Mo–Si–Zr ternary system [14]. The essence of this type of modelling is to derive the Gibbs energy functions (i.e. thermodynamic description) for phases in the Mo–Si–Zr system. From this thermodynamic description, thermodynamic properties and phase diagrams of the system can be calculated in a self-consistent manner. The thermodynamic description of the Mo–Si–Zr system was built upon the three constituent binaries, Mo–Si, Mo–Zr and Zr–Si. After thermodynamic descriptions for all the constituent binaries were established, the Gibbs energy of a ternary phase was calculated as the weighted average of the Gibbs energies of the binary phases using the Muggianu model [15], and additional contribution from ternary interactions. Details on the thermodynamic modelling of
the Mo–Si–Zr system and experimental validation can be found elsewhere [14]. Based on the thermodynamic description of Mo–Si–Zr, the Si solubility in the (Mo) phase was calculated for the two-phase ((Mo) and (Mo3Si)) alloys Mo–5Si and Mo–5Si–1Zr, as a function of temperature as shown in Fig. 6. The results indicates that Zr additions do not influence the Si solubility in the (Mo) phase significantly. 4. Discussion An important feature of MA is the formation of finegrained powders which allow control of the final microstructure and the strength level of the consolidated alloy. Mo–Si powders in which the domain size of a-Mo is significantly reduced during milling are no exception. The small domain size allows consolidation of the alloys at low temperatures due to accelerated diffusional processes through the grain boundaries and to obtain a compacted material with a grain size below 5 lm (Table 1). Since the strength in the low-temperature deformation mode depends on grain size according to the Hall–Petch relationship ry = r0 + kd 1/2, where r0 is the lattice friction stress for the dislocations, k the Hall–Petch constant, and d the grain size [16], it is obvious that the maximum strength at low temperatures (up to 816 °C) is higher than the value for
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Fig. 3. Fracture surfaces after deformation at room temperature of (a) Mo–1.5Si and (b) Mo–1.5Si–1Zr showing the positive effect of Zr addition on the grain boundary fracture strength.
Fig. 4. Fracture surfaces after deformation at room temperature of highpurity PM alloys (a) pure Mo and (b) Mo–0.3% Si showing the effect of minor Si additions on fracture mode and grain boundary fracture strength [5].
Mo–1.5Si with a grain size of 35 lm (600 MPa [5]). However, unlike the ductile behaviour of unalloyed Mo [5], the fine grained Mo–1.5Si solid solution fails within the elastic limit below 816 °C by intergranular fracture. The presence of impurities such as oxygen at the grain boundaries leads to embrittlement of unalloyed Mo [4]. This is also one potential reason for the brittle behaviour of the MAed Mo–1.5Si solid solutions which have oxygen concentrations in the range of 200–500 wt.ppm. However, as shown in Table 1 and Fig. 4 in the case of low oxygen (<5 wt.ppm) Mo(Si) solid solutions, concentrations as low as 0.3% Si lead already to brittle behaviour at room temperature. As shown by the Auger analyses, the grain boundaries in Mo(Si) solid solutions are enriched in Si. Even for bulk concentrations as low as 0.3% Si (see PM Mo–0.3Si) the grain boundaries appear to be saturated with Si. In view of the small size of Si atoms as compared to Mo, the pronounced segregation behaviour is not surprising. The absence of room temperature ductility in these alloys suggests that Si deteriorates the grain boundary strength, and therefore leads to a similar effect as oxygen. The lowering of the grain size by the addition of Y2O3 results in a significant improvement of the strength at low temperatures. However, no ductility was detected during
bend testing even though the submicron grain size would suggest plasticity [3]. Although no Auger analysis was performed for the Y2O3-containing alloy, the present results suggest that its grain boundaries were saturated with Si. In view of the small grain size and large Hall–Petch strengthening, the grain boundary strength was inadequate to allow plastic deformation at low temperatures. At 816 °C and above, however, the strength of the alloy is significantly reduced because diffusional creep processes become possible. An improvement of the ductility at low temperatures is obtained by the addition of Zr. Though having a larger grain size than the Y2O3-containing alloy, its bend strength is even higher because it does not fracture prior to the onset of macroscopic deformation. Further, the larger grain size of the Zr containing alloy leads to a higher strength level at temperatures above 538 °C as compared to the Y2O3-containing alloy (see Table 1). It is obvious that the grain boundary cohesion determines the ductility of the material. Plastic deformation is possible if the grain boundary cohesion is high enough that dislocation glide in the grain interior can be activated before the material fails intergranularly. Comparing the results of the Auger analyses of the Mo–1.5Si alloy with its Zr-containing counterpart, the only
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Fig. 5. Auger spectra of the grain interior which is identical for all alloys investigated (shown here for the MAed Mo–1.5Si alloy) and the grain boundary areas of PM Mo–0.3Si, MAed Mo–1.5Si, and MAed Mo–1.5Si–1Zr. Note the lower Si concentration at the grain boundaries of the Zr containing alloy when compared to the MAed Mo–1.5Si alloy.
obvious difference is the lower Si concentration at the grain boundaries of the latter. The lowering of the Si concentration is thought to increase the grain boundary cohesive strength and thus improve ductility. However, at this point it is not clear exactly how the Zr reduces the Si concentration at the grain boundaries. This could in principle happen by a reduction in the bulk solubility of Si due to Zr, but our thermodynamic calculations show that this does not happen. Therefore, it is likely that Zr (which is a much larger atom than Mo and therefore likely to segregate) competes with Si for the grain boundaries, thus reducing the Si concentration. It is also possible that the improvement of the grain boundary cohesive strength is not only due to the reduction in the Si concentration, but also due to the presence of Zr at the grain boundaries. Further experiments would be needed to sort this out. The results are, however, in qualitative agreement with earlier calculations by Seah [17], who showed in his simplified thermodynamics
Fig. 6. Solubility of Si in the (Mo) phase vs. temperature for alloy Mo– 5Si–1Zr and Mo–5Si.
H. Saage et al. / Acta Materialia 57 (2009) 3895–3901
approach that Zr additions to bcc metals such Fe or Mo are expected to be less detrimental to grain boundary embrittlement than Si or O. Since our analysis suggests that Zr competes with Si for grain boundary sites, the additions of Zr can indeed be considered ‘‘ductilizing” with respect to the binary Mo–Si material. Miller and Bryhan [9] found that oxygen is depleted at the grain boundaries of Mo alloys when elements such as Zr, C and B are segregated to the grain boundaries. Therefore, a further beneficial effect on the ductility may be due to the fact that Zr acts as getter for oxygen since it forms ZrO2 during consolidation (see Fig. 1). The ZrO2 formation lowers the oxygen concentration at the grain boundaries, and may therefore indirectly contribute to the grain boundary strength. Although the portion of transgranular fracture increases with the addition of Zr, a significant portion of intergranular fracture remains (see Fig. 4). This suggests that further improvement of the ductility by further strengthening of the grain boundaries, until the fracture mode becomes completely transgranular, is possible. This improvement might be reached by (i) further increase of the amount of Zr addition; (ii) further lowering the oxygen concentration by optimization of the sintering treatment; and (iii) further alloying additions which might lead to further depletion of Si and O at the grain boundaries. C, Ti and B are among the possible candidates for the latter two approaches [7,9,17,18]. 5. Conclusions 1. In Mo(Si) solid solutions Si segregates to the grain boundaries. The brittle behaviour of Mo(Si) at room temperature is, at least in part, caused by the reduction of the cohesive grain boundary strength due to segregated Si. 2. Addition of Y2O3 particles reduces the grain size and improves the room temperature strength, but plastic deformation is observed only at elevated temperatures. 3. Addition of Zr reduces not only the grain size, but results also in limited plastic deformation in room temperature bend tests. Strength levels of 2 GPa are obtained. The Zr getters detrimental oxygen by forming ZrO2 particles (which in turn help to pin the grain boundaries). In particular, Zr dramatically reduces the segregation of Si at the grain boundaries, thus enhancing their cohesive strength.
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4. Optimization of the beneficial effect of Zr and/or other alloying additions holds the promise of further improvements in the low-temperature ductility of Mo–Si alloys.
Acknowledgements The research was sponsored by the German Science Foundation (DFG) in the frame of the research unit 727 ‘‘Beyond Nickelbase Superalloys”. We are grateful to P. Je´hanno, M. Bo¨ning and H. Kestler for providing the material used in this study. J.H.S. acknowledges support of this work by the Division of Materials Sciences and Engineering, US Department of Energy, under Contract DE-AC05-00OR22725 with UT-Battelle, LLC, and support through the Schiebold Guest Professorship during a 5-month stay at the University of Magdeburg, Germany. References [1] Berczik DM. US patents 5,595,616 and 5,693,156; 1997. [2] Parthasarathy TA, Mendiratta M, Dimiduk DM. Acta Mater 2002;50:1857. [3] Northcott L. Molybdenum. New York: Academic Press; 1956. [4] Kumar A, Eyre BL. Proc Roy Soc London A 1980;370:431. [5] Sturm D, Heilmaier M, Schneibel JH, Je´hanno P, Skrotzki B, Saage H. Mater Sci Eng A 2007;463:107. [6] Geller CB, Smith RW, Hack JE, Saxe P, Wimmer E. Scripta Mater 2005;52:205. [7] Schneibel JH, Ritchie RO, Kruzic JJ, Tortorelli PF. Metall Mater Trans A 2005;36:525. [8] Miller MK, Bryhan AJ. Mater Sci Eng A 2002;327:80. [9] Perepezko JH, Sakidja R, Kim S. Mater. Res. Soc. Symp. Proc., vol. 646: Mater Res Soc, 2001, N4.5.1. [10] Suryanarayana C. Prog Mater Sci 2001;46:1. [11] Guha A. Bending strength tests. In: Metals handbook. Mechanical testing, vol. 8. American Society for Metals; 1985. p. 132. [12] Koch CC. Nanostruct Mater 1997;9:13. [13] Kru¨ger M, Franz S, Saage H, Heilmaier M, Schneibel JH, Je´hanno P, et al. Intermetallics 2008;16:933. [14] Yang Y, Schneibel JH. Thermodynamic modelling and mechanical testing for the multiphase Mo–Si–B–X (X = Ti, Zr, Hf) alloy for high temperature (1300 °C–1500 °C) services. Contract no. FA9550-08-C0024, Final report; 2008. [15] Muggianu YM, Gambino M, Bros JP. JCP 1975;72:83. [16] Haasen P. Physical Metallurgy. Cambridge: Cambridge University Press; 1978. [17] Seah MP. Acta Metall 1980;28:955. [18] Lutz H, Benesovsky F, Kiefer R. J Less-Common Metals 1968;16:249 [in German].
Acta Materialia 57 (2009) 3895–3901 www.elsevier.com/locate/actamat
Ductilization of Mo–Si solid solutions manufactured by powder metallurgy H. Saage a,*, M. Kru¨ger b, D. Sturm b, M. Heilmaier c, J.H. Schneibel b, E. George d, L. Heatherly d, Ch. Somsen e, G. Eggeler e, Y. Yang f b
a Faculty of Mechanical Engineering, University of Applied Sciences, Landshut, D-84036 Landshut, Germany Institut fu¨r Werkstoff- und Fu¨getechnik, Otto-von-Guericke-Universita¨t Magdeburg, D-39016 Magdeburg, Germany c Fachbereich Materialwissenschaft, TU Darmstadt, D-64287 Darmstadt, Germany d Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA e Institut fu¨r Werkstoffe, Ruhr-Universita¨t Bochum, D-44801 Bochum, Germany f Computherm LLC, 437 S. Yellow Stone Drive, Madison, WI 53719, USA
Received 18 March 2009; received in revised form 20 April 2009; accepted 23 April 2009 Available online 22 May 2009
Abstract Mo–1.5 at.% Si alloys with additions of either Y2O3 or Zr were manufactured by mechanical alloying. The Y2O3 particles reduced the grain size and increased the room temperature strength, but did not alleviate the brittleness of previously investigated Mo–1.5 at.% Si without Y2O3. Additions of Zr, on the other hand, resulted not only in a fine grain size and an extremely high bend strength (2 GPa), but also in limited bend ductility at room temperature. Zr additions are seen to be beneficial for three reasons. First, Zr reduces the grain size. Second, Zr getters detrimental oxygen by forming ZrO2 particles (which in turn help to pin the grain boundaries). Third, in situ Auger analysis shows that Zr reduces the concentration of Si segregated at the grain boundaries. This is thought to enhance the grain boundary cohesive strength and thus leads to the observed ductility. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mechanical alloying; Bending tests; Molybdenum silicides; Ductility; Microstructure
1. Introduction A family of three-phase Mo–Si–B alloys pioneered by Berczik [1] is considered as a potential new high-temperature material class for applications in air, since it combines the excellent high-temperature strength of Mo-base alloys (such as TZM: Mo–0.5Ti–0.08Zr–0.03C, wt.%) and the beneficial oxidation resistance of silicide phases. The two intermetallic phases Mo3Si and Mo5SiB2 (T2), which form in addition to the Mo(Si) solid solution (‘‘a-Mo”) matrix phase, improve the high-temperature oxidation resistance since they provide the Si and B necessary for forming a coating of a dense, protective SiO2(B) glass [2].
*
Corresponding author. Tel.: +49 871 506 219; fax: +49 871 506 9 219. E-mail address: [email protected] (H. Saage).
Mo possesses a body-centred cubic (bcc) crystal structure and is characterized by a brittle-to-ductile transition temperature (BDTT; see e.g. [3]) which is accompanied by a change from a mostly intergranular brittle fracture mode to mostly transgranular above the BDTT. The transition temperature itself depends on the purity of Mo, which tends to fail in a brittle manner by intergranular fracture at room temperature in the presence of even small amounts of interstitial impurities (50 wt.ppm) such as oxygen [3,4]. This is attributed to the weakening of the grain boundaries caused by this element. The solubility of Si in Mo depends on temperature but can be as high as 2 at.% at 1400 °C, the designated maximum application temperature of Mo–Si–B alloys. Therefore, in a previous work [5] we studied the mechanical behaviour of high-purity Mo(Si) solid solutions with Si contents ranging between 0.1 and 1 wt.% (i.e. about 0.3–3.3 at.%) as a reference for
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.04.040
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the matrix phase of three-phase Mo–Si–B alloys. As an important outcome we could demonstrate that while Si is one of the most effective solid solution strengtheners for Mo [3], it also strongly promotes the tendency towards brittle intergranular failure [5]. It is well known that the ductility of molybdenum can be improved by reducing its grain size [3]. Thus, one of the approaches to ductilize Mo(Si) solid solutions in the present work is to add Y2O3, a thermodynamically very stable oxide, which reduces grain boundary mobility, and therefore stabilizes a small grain size during powder metallurgical (PM) manufacturing. However, this may have a detrimental impact on the creep performance at higher temperatures at which the material is intended to be used. Another approach is to strengthen the grain boundaries by reducing the oxygen concentration at the grain boundaries by adding carbon [4]. Therefore, a second and possibly more viable way is to find suitable microalloying additions which improve the ductility by reducing the tendency of solute atoms to segregate to the grain boundaries and embrittle them. Among the possible candidates which have been identified through a theoretical approach [6], Zr has been shown recently to improve the fracture toughness of Mo–12Si– 8.5B (at.%) [7]. According to Miller and Bryhan [8], who measured the distribution of elements at and close to the grain boundaries by atom probe tomography, the basic mechanism for the improvement of the grain boundary fracture stress of a welded Mo-base alloy containing small amounts of Zr, C and B was found to be the substitution of O and possibly N by these additions at the grain boundaries. Therefore, in our second approach we use ZrH2, which decomposes into its elements during manufacturing at high temperatures, in order to ductilize Mo(Si) solid solutions prepared by PM. 2. Materials and methods The maximum solubility of Si in Mo at the targeted application temperature of at least 1200 °C is about 1.5 at.% (estimated from extrapolation of the respective ternary phase diagram) [9]. Thus, elemental powders of Mo and Si of 99.95% and 99.9% purity, respectively, were mixed to obtain this target composition (compositions are given in at.% unless stated otherwise). About 0.7 vol.% of Y2O3 powder or about 1 wt.% ZrH2 powder, which is reduced to about 1 at.% Zr upon sintering, were added to obtain two additional compositions. Mechanical alloying (MA, a high-energy milling process [10]) was carried out for 20 h under a protective (argon) atmosphere in a planetary ball mill with a rotational speed of 200 rpm and a powder to ball weight ratio of 1:13. In the following all alloy designations without a prefix refer to alloys where MA was used to prepare the powders. In some cases PM is used as prefix instead which refers to alloys prepared from blended powders (not MAed). The detailed processing of these alloys is described elsewhere [5]. After milling
the powders were cold isostatically pressed at 200 MPa and sintered in hydrogen at 1500 °C to reduce the impurity content (mainly O, N, H). Depending on the preparation technique the powders respond differently to this treatment. Simple mixtures of the elements Mo and Si have impurity levels of below 5 wt.ppm [5], whereas the MAed powders studied here possess 200–500 wt.ppm oxygen after sintering, which remains at this level upon further consolidation. As a final step the sintered specimens were further consolidated by hot isostatic pressing (HIP) at 1500 °C and 200 MPa for 4 h. Electrodischarge machining was used to prepare bending specimens with a rectangular cross-section of about 3.8 3 mm and a length of 45 mm from the HIP billets and final surface quality was attained by grinding. The mechanical properties were assessed by three-point bending tests with a span of 40 mm in a temperature range of room temperature–1093 °C using a Zwick electromechanical testing machine equipped with a Maytec furnace at a crosshead speed of 0.01 mm min 1. At temperatures above 538 °C a gas mixture of argon with 2% hydrogen was used for the bend tests to prevent oxidation. Bend deformation was continuously monitored by Al2O3 rods attached to the centre of the lower (tensile) side of the specimens and to the lower pushrod, respectively. The stresses and (elastic) strains on the tensile side of the bend specimen were calculated utilizing elastic beam theory [11], which is sufficient to detect the BDTT, i.e. the temperature above which plasticity occurs. With this configuration even small deviations from the elastic behaviour could clearly be resolved. Specimens with the composition Mo–1.5Si and Mo– 1.5Si–1Zr were broken in situ in the ultra-high vacuum of an Auger microscope (PHI 680 SAM) and their fracture surfaces were analysed for possible segregation of the alloying elements at the grain boundaries. From secondary electron (SE) scanning electron microscopy images of the fracture surfaces, median values of the grain size were determined by directly measuring at least 250 grains. 3. Results 3.1. Mechanical alloying and consolidation Mechanical alloying of the Mo–1.5Si powder mixtures is accompanied by a significant microstructural refinement. ˚ durThe Mo domain size is reduced to values below 20 A ing milling and, typical for brittle–ductile-type powder mixtures [10,12], the additional phases Y2O3 and ZrH2 are first crushed into smaller pieces and subsequently introduced into the more ductile Mo powder particles to form fine, homogeneously distributed dispersions. Due to the substitution of Mo by smaller Si atoms the lattice parameter is reduced during MA and the dissolution of 1.5% Si is complete after milling for 20 h [5,13]. The ZrH2 decomposes during consolidation by releasing the hydrogen. Some of the Zr is thought to remain in solid solution, while the
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remainder reacts with residual oxygen to form ZrO2 particles in the grain interior and at the grain boundaries (Fig. 1). The Si remains in solution after consolidation (HIPing) at 1500 °C. As shown in Table 1 the grain size after consolidation depends on both the composition as well as the processing conditions. While the median value of the grain size of the Mo(Si) solid solution obtained from MAed powder is about 3.2 lm, the addition of 1% Zr reduces the grain size to 1.7 lm. Alloying with Y2O3 dispersoids led to the lowest grain size, less than 0.5 lm, after consolidation. For comparison, specimens consolidated from simply mixed Mo– Si powders of our previous study [5] yielded mean grain sizes which were larger by more than one order of magnitude.
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tional Y2O3 dispersoids are brittle at temperatures up to 538 °C. However, both alloys show some plastic deformability at 816 °C. In contrast, the alloy with ZrH2 addition already reveals limited plastic strain at room temperature at a very high strength level of around 2 GPa. Moreover, this alloy reaches a plastic strain of about 1.4% at 538 °C while exceeding the strength level of 2 GPa. As expected, at even higher temperatures the strength of this alloy decreases while the plastic deformability increases. Fracture of the Mo(Si) solid solution occurs predominantly along the grain boundaries with a fraction of intergranular failure of about 93%, with the remainder showing transgranular fracture. In contrast, the Zr-containing alloy shows less intergranular fracture (80%, see Fig. 3), whereas fracture of the Y2O3-containing alloy occurs almost 100% along the grain boundaries at room temperature. For comparison, on the basis of tensile tests, we have shown [5] that the addition of only about 0.3 % Si leads already to brittle intergranular fracture at room temperature, while nominally pure Mo almost reaches 15% plastic strain before failure and reveals a transcrystalline fracture mode (see Table 1 and Fig. 4).
3.2. Temperature dependence of strength and ductility Stress–strain (calculated for the outer tensile fibre) curves for the alloys tested at room temperature and at 538 °C are compared in Fig. 2. Values for the yield strength, maximum bend strength and plastic strain before failure for all test temperatures are given in Table 1. The MAed Mo–1.5 Si solid solution and the alloy with addi-
3.3. Auger analysis on the fracture surface
Fig. 1. Bright-field TEM image (two beam condition) showing fine dispersed ZrO2 particles in the grain interior and the pinning of dislocations.
In Fig. 5a Auger spectra of the grain boundary regions (intergranular fracture areas) and the grain interior (transgranular) are compared for PM Mo–0.3Si as well as for MA Mo–1.5Si and MA Mo–1.5Si–1Zr solid solution. Unfortunately, it was not possible to reliably compare the oxygen concentration at the grain boundaries. Also, because of peak overlap with Mo, the Zr concentration could not be determined. The only obvious difference of the spectra taken from the grain boundaries and the grain interior is in the size of the Si peak at 91 eV. While this peak is pronounced for the grain boundaries, the transgranular fracture area (grain interior) shows, consistent with the low bulk Si concentration, only a small peak. Quantitative analyses of the Si concentration on the basis of several Auger spectra shown in Fig. 5b confirm that the Si concentration is reduced at the grain boundaries of the Zr containing Mo–1.5Si solid solution, as compared to Mo(Si) alloys without Zr. Hence, Si is strongly segregated at the grain boundaries in all binary Mo–Si alloys
Table 1 Different parameters of the MAed materials investigated here. In addition, tensile test values for two PM Mo–Si solid solutions which were taken from Ref. [5] are also shown. Alloy
Median of the Yield strength in MPa grain size in lm
Maximum bend or tensile strength Plastic strain in % in MPa
RT 538 °C 816 °C 1093 °C RT
538 °C
816 °C
1093 °C
RT 538 °C 816 °C 1093 °C
Mo–1.5Si (MAed) 3.2 Mo–1.5Si + 0.7 vol.% Y2O3 (MAed) 0.45 1.7 Mo–1.5Si + 1 wt.% ZrH2 (MAed)
– – –
– – 1950
920 960 1040
– 67 (3%) 385
– – – – 0.1 1.4
0.3 1.2 0.9
– >3% >3%
Mo–0.3Si (PM) Mo–1.5Si (PM)
– –
166 –
150 331
219 296
– –
27 3.3
90 0.6
54 35
– 48 305
516 1581 2001
731 1804 2075
953 991 1203
132 275
– –
377 –
365 541
32 –
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Fig. 2. Stress–strain curves (outer fibre) of pure Mo–1.5Si and Mo–1.5Si with Y2O3 and ZrH2 addition at (a) room temperature and at (b) 538 °C, (c) 816 °C and (d) 1093 °C.
but to a much lesser extent in the Zr-containing alloy, indicating that the presence of Zr leads to depletion of Si at the grain boundaries. 3.4. Calphad-type thermodynamic modelling of the Mo–Si– Zr ternary system The question arises whether the reduction of the Si grain boundary concentration due to Zr could be caused by a reduction in the bulk solubility in the presence of Zr. Therefore, the effect of Zr addition on the solubility of Si in Mo was studied by Calphad-type thermodynamic modelling of the Mo–Si–Zr ternary system [14]. The essence of this type of modelling is to derive the Gibbs energy functions (i.e. thermodynamic description) for phases in the Mo–Si–Zr system. From this thermodynamic description, thermodynamic properties and phase diagrams of the system can be calculated in a self-consistent manner. The thermodynamic description of the Mo–Si–Zr system was built upon the three constituent binaries, Mo–Si, Mo–Zr and Zr–Si. After thermodynamic descriptions for all the constituent binaries were established, the Gibbs energy of a ternary phase was calculated as the weighted average of the Gibbs energies of the binary phases using the Muggianu model [15], and additional contribution from ternary interactions. Details on the thermodynamic modelling of
the Mo–Si–Zr system and experimental validation can be found elsewhere [14]. Based on the thermodynamic description of Mo–Si–Zr, the Si solubility in the (Mo) phase was calculated for the two-phase ((Mo) and (Mo3Si)) alloys Mo–5Si and Mo–5Si–1Zr, as a function of temperature as shown in Fig. 6. The results indicates that Zr additions do not influence the Si solubility in the (Mo) phase significantly. 4. Discussion An important feature of MA is the formation of finegrained powders which allow control of the final microstructure and the strength level of the consolidated alloy. Mo–Si powders in which the domain size of a-Mo is significantly reduced during milling are no exception. The small domain size allows consolidation of the alloys at low temperatures due to accelerated diffusional processes through the grain boundaries and to obtain a compacted material with a grain size below 5 lm (Table 1). Since the strength in the low-temperature deformation mode depends on grain size according to the Hall–Petch relationship ry = r0 + kd 1/2, where r0 is the lattice friction stress for the dislocations, k the Hall–Petch constant, and d the grain size [16], it is obvious that the maximum strength at low temperatures (up to 816 °C) is higher than the value for
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Fig. 3. Fracture surfaces after deformation at room temperature of (a) Mo–1.5Si and (b) Mo–1.5Si–1Zr showing the positive effect of Zr addition on the grain boundary fracture strength.
Fig. 4. Fracture surfaces after deformation at room temperature of highpurity PM alloys (a) pure Mo and (b) Mo–0.3% Si showing the effect of minor Si additions on fracture mode and grain boundary fracture strength [5].
Mo–1.5Si with a grain size of 35 lm (600 MPa [5]). However, unlike the ductile behaviour of unalloyed Mo [5], the fine grained Mo–1.5Si solid solution fails within the elastic limit below 816 °C by intergranular fracture. The presence of impurities such as oxygen at the grain boundaries leads to embrittlement of unalloyed Mo [4]. This is also one potential reason for the brittle behaviour of the MAed Mo–1.5Si solid solutions which have oxygen concentrations in the range of 200–500 wt.ppm. However, as shown in Table 1 and Fig. 4 in the case of low oxygen (<5 wt.ppm) Mo(Si) solid solutions, concentrations as low as 0.3% Si lead already to brittle behaviour at room temperature. As shown by the Auger analyses, the grain boundaries in Mo(Si) solid solutions are enriched in Si. Even for bulk concentrations as low as 0.3% Si (see PM Mo–0.3Si) the grain boundaries appear to be saturated with Si. In view of the small size of Si atoms as compared to Mo, the pronounced segregation behaviour is not surprising. The absence of room temperature ductility in these alloys suggests that Si deteriorates the grain boundary strength, and therefore leads to a similar effect as oxygen. The lowering of the grain size by the addition of Y2O3 results in a significant improvement of the strength at low temperatures. However, no ductility was detected during
bend testing even though the submicron grain size would suggest plasticity [3]. Although no Auger analysis was performed for the Y2O3-containing alloy, the present results suggest that its grain boundaries were saturated with Si. In view of the small grain size and large Hall–Petch strengthening, the grain boundary strength was inadequate to allow plastic deformation at low temperatures. At 816 °C and above, however, the strength of the alloy is significantly reduced because diffusional creep processes become possible. An improvement of the ductility at low temperatures is obtained by the addition of Zr. Though having a larger grain size than the Y2O3-containing alloy, its bend strength is even higher because it does not fracture prior to the onset of macroscopic deformation. Further, the larger grain size of the Zr containing alloy leads to a higher strength level at temperatures above 538 °C as compared to the Y2O3-containing alloy (see Table 1). It is obvious that the grain boundary cohesion determines the ductility of the material. Plastic deformation is possible if the grain boundary cohesion is high enough that dislocation glide in the grain interior can be activated before the material fails intergranularly. Comparing the results of the Auger analyses of the Mo–1.5Si alloy with its Zr-containing counterpart, the only
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Fig. 5. Auger spectra of the grain interior which is identical for all alloys investigated (shown here for the MAed Mo–1.5Si alloy) and the grain boundary areas of PM Mo–0.3Si, MAed Mo–1.5Si, and MAed Mo–1.5Si–1Zr. Note the lower Si concentration at the grain boundaries of the Zr containing alloy when compared to the MAed Mo–1.5Si alloy.
obvious difference is the lower Si concentration at the grain boundaries of the latter. The lowering of the Si concentration is thought to increase the grain boundary cohesive strength and thus improve ductility. However, at this point it is not clear exactly how the Zr reduces the Si concentration at the grain boundaries. This could in principle happen by a reduction in the bulk solubility of Si due to Zr, but our thermodynamic calculations show that this does not happen. Therefore, it is likely that Zr (which is a much larger atom than Mo and therefore likely to segregate) competes with Si for the grain boundaries, thus reducing the Si concentration. It is also possible that the improvement of the grain boundary cohesive strength is not only due to the reduction in the Si concentration, but also due to the presence of Zr at the grain boundaries. Further experiments would be needed to sort this out. The results are, however, in qualitative agreement with earlier calculations by Seah [17], who showed in his simplified thermodynamics
Fig. 6. Solubility of Si in the (Mo) phase vs. temperature for alloy Mo– 5Si–1Zr and Mo–5Si.
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approach that Zr additions to bcc metals such Fe or Mo are expected to be less detrimental to grain boundary embrittlement than Si or O. Since our analysis suggests that Zr competes with Si for grain boundary sites, the additions of Zr can indeed be considered ‘‘ductilizing” with respect to the binary Mo–Si material. Miller and Bryhan [9] found that oxygen is depleted at the grain boundaries of Mo alloys when elements such as Zr, C and B are segregated to the grain boundaries. Therefore, a further beneficial effect on the ductility may be due to the fact that Zr acts as getter for oxygen since it forms ZrO2 during consolidation (see Fig. 1). The ZrO2 formation lowers the oxygen concentration at the grain boundaries, and may therefore indirectly contribute to the grain boundary strength. Although the portion of transgranular fracture increases with the addition of Zr, a significant portion of intergranular fracture remains (see Fig. 4). This suggests that further improvement of the ductility by further strengthening of the grain boundaries, until the fracture mode becomes completely transgranular, is possible. This improvement might be reached by (i) further increase of the amount of Zr addition; (ii) further lowering the oxygen concentration by optimization of the sintering treatment; and (iii) further alloying additions which might lead to further depletion of Si and O at the grain boundaries. C, Ti and B are among the possible candidates for the latter two approaches [7,9,17,18]. 5. Conclusions 1. In Mo(Si) solid solutions Si segregates to the grain boundaries. The brittle behaviour of Mo(Si) at room temperature is, at least in part, caused by the reduction of the cohesive grain boundary strength due to segregated Si. 2. Addition of Y2O3 particles reduces the grain size and improves the room temperature strength, but plastic deformation is observed only at elevated temperatures. 3. Addition of Zr reduces not only the grain size, but results also in limited plastic deformation in room temperature bend tests. Strength levels of 2 GPa are obtained. The Zr getters detrimental oxygen by forming ZrO2 particles (which in turn help to pin the grain boundaries). In particular, Zr dramatically reduces the segregation of Si at the grain boundaries, thus enhancing their cohesive strength.
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4. Optimization of the beneficial effect of Zr and/or other alloying additions holds the promise of further improvements in the low-temperature ductility of Mo–Si alloys.
Acknowledgements The research was sponsored by the German Science Foundation (DFG) in the frame of the research unit 727 ‘‘Beyond Nickelbase Superalloys”. We are grateful to P. Je´hanno, M. Bo¨ning and H. Kestler for providing the material used in this study. J.H.S. acknowledges support of this work by the Division of Materials Sciences and Engineering, US Department of Energy, under Contract DE-AC05-00OR22725 with UT-Battelle, LLC, and support through the Schiebold Guest Professorship during a 5-month stay at the University of Magdeburg, Germany. References [1] Berczik DM. US patents 5,595,616 and 5,693,156; 1997. [2] Parthasarathy TA, Mendiratta M, Dimiduk DM. Acta Mater 2002;50:1857. [3] Northcott L. Molybdenum. New York: Academic Press; 1956. [4] Kumar A, Eyre BL. Proc Roy Soc London A 1980;370:431. [5] Sturm D, Heilmaier M, Schneibel JH, Je´hanno P, Skrotzki B, Saage H. Mater Sci Eng A 2007;463:107. [6] Geller CB, Smith RW, Hack JE, Saxe P, Wimmer E. Scripta Mater 2005;52:205. [7] Schneibel JH, Ritchie RO, Kruzic JJ, Tortorelli PF. Metall Mater Trans A 2005;36:525. [8] Miller MK, Bryhan AJ. Mater Sci Eng A 2002;327:80. [9] Perepezko JH, Sakidja R, Kim S. Mater. Res. Soc. Symp. Proc., vol. 646: Mater Res Soc, 2001, N4.5.1. [10] Suryanarayana C. Prog Mater Sci 2001;46:1. [11] Guha A. Bending strength tests. In: Metals handbook. Mechanical testing, vol. 8. American Society for Metals; 1985. p. 132. [12] Koch CC. Nanostruct Mater 1997;9:13. [13] Kru¨ger M, Franz S, Saage H, Heilmaier M, Schneibel JH, Je´hanno P, et al. Intermetallics 2008;16:933. [14] Yang Y, Schneibel JH. Thermodynamic modelling and mechanical testing for the multiphase Mo–Si–B–X (X = Ti, Zr, Hf) alloy for high temperature (1300 °C–1500 °C) services. Contract no. FA9550-08-C0024, Final report; 2008. [15] Muggianu YM, Gambino M, Bros JP. JCP 1975;72:83. [16] Haasen P. Physical Metallurgy. Cambridge: Cambridge University Press; 1978. [17] Seah MP. Acta Metall 1980;28:955. [18] Lutz H, Benesovsky F, Kiefer R. J Less-Common Metals 1968;16:249 [in German].