Mechanical properties of molybdenum disilicide based materials consolidated by hot isostatic pressing (HIP)

Acta metall, mater. Vol. 42, No. 11, pp. 3751 3757, 1994

Pergamon

0956-7151(94)E0149-B

Copyright © 1994ElsevierScienceLtd Printed in Great Britain.All rights reserved 0956-7151/94$7.00+ 0.00

MECHANICAL PROPERTIES OF MOLYBDENUM DISILICIDE BASED MATERIALS CONSOLIDATED BY HOT ISOSTATIC PRESSING (HIP) R. SURYANARAYANAN, S. M. L. SASTRY and K. L. 3ERINA Materials Science and Engineering Program, Department of Mechanical Engineering, Washington University, 1 Brookings Drive, St Louis, MO63130, U.S.A. (Received 28 February 1994)

Abstract--The influence of hot isostatic pressing (HIP) consolidation parameters on the mechanical properties of molybdenum disilicide (MoSi2) and MoSi2 reinforced with ductile and brittle reinforcements was studied. MoSi2, MoSi2 20 vol.% coarse and fine niobium powder and MoSi2-20vol.% siliconcarbide whiskers consolidated by HIP at 1200-1400°C, 207 MPa, for l and 4 h were tested in compression for elevated temperature strength and creep resistance. Single-edge-notchedspecimens of the three materials were tested in a three-point bend configuration for fracture toughness. Mechanical properties were related with consolidation parameters and post-HIP microstructures.

1. INTRODUCTION Applications in advanced aerospace structural cornponents require a balance of properties such as low density, creep resistance and strength at elevated temperatures, damage tolerance, and environmental resistance. The silicide compound, MoSi 2, has been a popular candidate material system for such applications [1]. The strength and ductility [2-8], creep [9-13], fracture toughness [2, 4, 13-19], hardness [14, 20], and fatigue [2] properties of fully dense MoSi 2 and MoSi 2 reinforced with niobium, silicon carbide, etc., have been studied extensively in the recent past. The results have been reviewed by Vasudevan and Petrovic [21] and in a more general format by Kumar and Liu [22]. Several issues, however, are still unresolved. The ductile-to-brittle transition temperature (DBTT), the mechanisms of creep, effects of type and volume fraction of reinforcements on creep properties, the micromechanisms of fracture and the porosity dependence of fracture properties are some of the topics which need further investigation, Mechanical properties of monolithic MoSi 2 have been a major hindrance for their use in structural applications. With a ductile-to-brittle transition ternperature in the range of 1000°C [23] to 1300°C [24], it has low toughness at low temperatures and low strength above the DBTT. Two possible methods to overcome both these problems are alloying or the addition of suitable reinforcements. Improvements in hardness properties can be achieved by additions of small amounts of carbon, as observed by Maloy et al. [20]. Studies on single crystal deformation of MoSi2

conducted by Umakoshi et al. [25], Hirano et al. [26], and Kimura et al. [27] indicate that crystallographic texture control by deformation processing can improve strength. An important requirement for high temperature applications is satisfactory creep resistance. The creep properties of MoSi2 have been successfully enhanced by the addition of SiC whiskers [10]. Although creep mechanisms in MoSi 2 composites are not clearly understood, it has been speculated that the overall process is due to the combined effects of grain boundary sliding and dislocation glide/climb at high applied stresses, while grain boundary sliding domihates at lower stresses [9, 12]. Second phase additions of brittle or ductile reinforcements to MoSi 2 result in improvements in room temperature fracture toughness and high temperature strength and creep resistance, with no deleterious effect on oxidation resistance. Reinforcements like SiC, ZrO2, TiB2, and Nb improve room temperature fracture toughness of MoSi2. The objective of this work was to characterize mechanical properties of MoSi 2 based samples consolidated under different HIP conditions. Details of the consolidation of samples by HIP are given elsewhere [28]. Elevated temperature strength and creep resistance, and room temperature fracture toughness properties were investigated. Mechanical properties were related to HIP consolidation parameters and post-HIP microstructural features. Tests were done at conditions similar to HIP process parameters in order to experimentally measure material property data (e.g. creep data) for use in HIP modeling and construction of HIP diagrams [28].

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SURYANARAYANAN eta/.:

MECHANICAL PROPERTIES OF MoSi2

2. EXPERIMENTAL METHODS

The apparent stress intensity factor, KQ, for the bend test specimen was calculated using the following equation [31]

2.1. Compression strength Compacts of MoSi 2, MoSi2-Nb, and MoSi2-SiC produced for the consolidation experiments [28] were machined into rectangular pieces by electro discharge machining. Specimens tested were the highest density

Po" S K° = B . W 3.2'f ( a / W ) where 1

samples of MoSi2, MoSi2-Nb, and MoSi2-SiC. Specimens wereapproximately 10mm x 6 m i n x 3ram in size. Specimens were soaked at the test temperatures of 1100 and 1300°C before loading. Compression tests were performed on a MTS closed loop servohydraulic test machine at an initial strain rate of 5 x 10 -4 s -1.

2.2. Creep resistance Specimens were tested under constant load compression, in the temperature range of 1100-1300°C, using SiC platens and push rods. The displacement of the lower cross head of the MTS testing machine was recorded as a function of time. The incremental load method was used once steady state creep was attained at a particular load.

2.3. Fracture toughness A single edge notch was introduced in the middle of each specimen using a 0.3 mm diamond cutting wheel. The rectangular specimens were approx, 25 x 6 × 3 mm and the initial crack lengths produced by the diamond wheel were approx. 2.5 mm. Specimens were loaded monotonically to failure under three point bending. Fracture toughness measurements of fatigue pre-cracked and non-precracked specimens of MoSi2 based materials indicated no significant differences [29]. No fatigue pre-cracking was done on any of the samples tested. Munz et al. [30] have shown significant effect of the notch root radius on the measured fracture toughness of sintered aluminum oxide. However, if the notch is fine enough (66 p m for A1203) , the machined notch acts as a sharp crack. Since the notch root radius in the present case was ~ 150 pm, measured toughness values are not necessarily equal to the plane strain fracture toughness values of these materials. We have used KQ instead of K~c as the toughness parameter although we believe the difference between the two values is not significant. It must be pointed out that the principal objective of this part of the work is to correlate toughness with relative density and compare the behavior of the three materials in the study, and not specifically the precise determination of K~c. Reasonable scatter in data and consistency of results in all cases examined suggests validity of the conclusions we arrive at in later sections,

(1)

W/\

W

where: PQ = load in kN (klbf); B = specimen thickness in cm (inches); S = s p a n in cm (inches); W = specimen width in cm (inches); a = crack length in cm (inches). Fracture surface morphologies of the most dense and the least dense samples in each type of material were studied by a H I T A C H I $4500 Scanning Electron Microscope. Fracture modes in each case were identified. 3. RESULTS AND DISCUSSION

3.1. Compression strength The compressive yield strengths of the highest density samples of MoSi 2, MoSi2-Nb, and MoSi2-SiC, along with results of earlier studies are shown in Table 1. Compression strength data for the monolithic case shows that attractive strength is retained at least until 1300°C. Measured strength at 1100°C is much higher than data from literature for similar temperatures. This may be due to lower levels of the silica rich grain boundary phase in the present case than in earlier studies where consolidation was carried out at much higher temperatures. Optimal processing, therefore, may be a means of reducing the undesirable silica phase in the matrix. Tested samples were intact after compressive deformation of ~ 2 . 5 % strain. This ductile squashing indicates that the DBTT is at a temperature less than 1100°C. This observation, although conforming to the generally accepted value of DBTT, is in contrast to the findings of Aikin [24] and Schwartz [29]. The DBTT of MoSi 2, therefore, is an unresolved question and needs further investigation. In the case of the Nb-reinforced composite, strengths at 1100 and 1300°C are comparable to values in previous studies at similar temperatures. Samples showed buckling at the later stages of compression at both test temperatures. The strengths are lower in MoSi2-SiC specimens, compared to the other two cases. Compression strengths of the present work are only comparable with the four-point flexural strengths of Refs [5, 8]. The poor strengths at these temperatures may be attributed to the larger levels of the grain boundary Si-rich phase that may be viscous at these temperatures [22]. As suggested by Maloy et al. [20] a small

SURYANARAYANAN et al.: MECHANICAL PROPERTIES OF MoSi2

3753

Table 1. Effect of temperatureon the compressionstrength (in MPa) of MoSi2based materials MonolithicMoSi2 T ("C) 1050 1100 1150 1200 1250 1300 1350 1400 ~ (s ~) Ref. [7l 594 493 305 222 5 x 10-4 Ref. [8l 325 194 7 x 10 5 Present work 826.7 394.4 5 x 10 4

T (°C)

1050

Ref. [7] Present work

513

T (°C) Refl [9], flexure Ref. [10], flexure Present work

MoSi2-Nb 1150 1250

1100

392

1300

286

156

450.2

1100

265.5

MoSi2-SiC 1200 1300 236 394

295.8

1350

1400 39 73

171

addition of carbon to MoSi2 can reduce this problem since carbon acts as a deoxidant, thereby removing the silica grain boundary phase. Also C reacts with MoSi2 to form SiC leading to increase in toughness, Samples crumbled after deforming in compression to ,-~2.5% strain, indicating that there was no or minimal plastic deformation in this case. A higher level of porosity ( > 4%) might also be a controlling factor for the brittleness of the SiC reinforced material. Optical micrographs of the SiC reinforced material showed agglomeration of the SiC whiskers, possibly due to electrostatic attractions [28]. This agglomeration of the SiC whiskers causes regions of the composite to be highly brittle compared to the matrix. We believe this is a major reason for the low strengths in this material. A better distribution of the whiskers in the matrix may lead to a significantly higher strength value than measured in the present study, The scatter in strength data for MoSi2-SiC has been investigated by Ting [32]. In his article, a Weibull analysis of 45 samples of MoSi2-SiC tested for flexural strength showed a wide range of values. A Weibull modulus of 5 was calculated for the experimental data on strength for this material. This is an indicator of the wide range of values of strength that are expected for MoSi 2 based materials. It is clear that a detailed investigation is necessary to arrive at some confident values of the strength of these materials.

for MoSi2-SiC in earlier studies [9-13, 21, 33]. One possible explanation for this discrepancy is the low density of the MoSi2-SiC specimen, approx. 96%, compared to the nearly theoretical dense specimens used by earlier workers. Also, the much smaller grain sizes in the SiC reinforced case may enhance the kinetics of Mo diffusion in MoSi2, which being the rate controlling step in the kinetics of creep of MoSi2 [9] may lead to higher creep rates. French et al. [34] have shown that for significant increase in creep resistance SiC volume fractions need to be in excess of 40%. The average creep stress exponents for MoSi2, MoSi2-Nb, and MoSi2-SiC in the 1200-1300°C temperature range are 2.2, 2.55 and 2.41, respectively. The stress exponent for the Nb reinforced material varies from 2.06 at 1300°C to 4.02 at 1100°C. This increase in the value of n with a decrease in temperature is an expected result for these materials. The plot of steady state creep rate vs reciprocal absolute temperature was used to determine the activation energy for the MoSi2-Nb material as 160 kJ/mol [35].

3.3. Fracture toughness 3.3. I. Monolithic MoSi2. Fracture in MoSi 2 occurs predominantly by cleavage at temperatures below Stress(MPa) 10o -4.0 . . . . . .

loo0 , ........

104

2 3.2. Creep resistance Figure 1 shows the steady state creep rate as a function of applied stress for the three materials at different temperatures. The slopes of the regression fit lines [corresponding to the power law creep (PLC) stress exponents] are noted in the figure. The stress exponents are comparable to results of earlier studies on MoSi 2 and MoSi2-SiC [9-13, 21, 33]. The creep rates of MoSi2-Nb are higher at both temperatures compared to the monolithic case. The creep rates of MoSi2-SiC is slightly lower than the monolithic material. This small difference in creep rates is contrary to the significantly lower creep rates obtained

-4.5 ~ ~ -5.o -s.s _g' ~" -6.0

.~. / ~

~'~# o lo-~

a~ ~ .%/ v

f

ouosi,,~2oooc °uost~~3°°°e

:~

~, MoSi=-Nb,1200°C // eMoS~-Nb, 1300°C /=MoSt.Nb. 1100oc

10" / • MoSi2-SIC,13000C / OMoSi2.SiC,1200*C "6"57.7 ' 7'.9' 8'.1 '8.3 ' '5' 8'.7 ' 819 ' 911 ' 9.3 Iogor(pa)

Fig. 1. Steady state creep rate as a function of applied stress for MoSi2, MoSi2-Nb and MoSiz-SiC.

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SURYANARAYANAN et al.: MECHANICAL PROPERTIES OF MoSt 2

DBTT, cleavage plus intergranular fracture in the 1200-1400°C range, and by transgranular fracture at higher temperatures [2, 18]. A detailed study of fracture modes in MoSt2, by Wade and Petrovic [18], concluded that transgranular fracture is the primary fracture mode in this material. The avera~ge fracture toughness was found to be 3 MPax/m. It was suggested by Wade and Petrovic [18] that the high percentage of transgranular fracture implies that crystal cleavage planes are weaker than grain boundaries. The transgranular fracture mode may be attributed to the crystallographic anisotropy and layered structure of MoSt 2. The large c / a ratio (2.45) may be expected to generate anisotropic stresses within individual grains after processing, contributing to transgranular fracture behavior. A combination of the above mentioned factors may adequately explain the brittle transgranular fracture of MoSt: at room temperature [18]. Figure 2 shows the results of the present experimental work as a toughness vs relative density plot for monolithic MoSt2. For relative density less than 0.95, KQ is constant at ~ 3 MPax/-~ and increases to ~ 3.3 M P a x / ~ and ~4.0 M P a v / m for powders consolidated for 1 and 4 h, respectively. The toughness value increase above 95% density corresponds to the consolidation stage where pores become isolated and more spherical. This increase in toughness is not surprising since in the stage when pores are interconnected, cracks have existing channels to propagate through, leading to quicker failure, Fracture surface morphologies of the MoSt 2 samples consolidated at 1200 and 1400°C are shown in Fig. 3(a,b). The sample consolidated at 1200°C shows a cleavage fracture mode. The cusp shaped pores (indicated by arrows) are also seen which are indicative of a power law creep densification mechanism. For the sample consolidated at 1400°C, smaller grains fracture by an intergranular mode and larger grains show cleavage fracture. This is because of the duplex grain structures observed in this material [28]. 6.0

. o

.

.

.

.

.

.

.

o"

MoSiz, 1hr.

o D

• MoSi2, 4 hrs.

UoSia-Nb(C),1hr. zxMoSJ,.SiC,1hr.

[]

t~

4.5

• MoSi=-SiC, 4hrs.

c] cJ

• • 8

E ~.

3.0

o

•

°

zx 1.5

•

•I t 0.90 0.95 t.00 Relative Density, p Fig. 2. Toughness as a function of relative density for monolithic MoSi:, MoSi2-Nb (coarse) and MoSi2-SiC. 0"0.75

I

I

0.80

0.85

3.3.2. M o S i 2 - 2 O v o l . % Nb. Ductile reinforcements have been used effectively to increase fracture toughness of brittle matrices [36-41]. Previous studies have utilized niobium [2, 15-17, 29, 42], tantalum [19], and T i - N b [2, 43], as ductile reinforcements in a MoSt 2 matrix. However, the resulting composites are unstable at elevated temperatures because, for example, niobium reacts with MoSt 2 forming brittle intermetallics at the interface. The formation of these cornpounds degrade the toughening effect, and some investigators have studied the effectiveness of inert diffusion barrier coatings like A1203, ZrO2 [17], and Y203 [15] to counter this problem. As an alternative to the coating approach, the methodology of the present investigation was to reduce the HIP temperature to a minimum without compromising on the level of density, thereby, minimizing the reaction products at the interface. A study of fracture toughness of MoSt 2 reinforced with 20 vol.% Nb [2] found that the ductile phase promotes a modest 24-33% increase in fracture toughness when toughness is expressed in terms of the stress intensity factor, K. Toughening due to ductile reinforcements has been generally attributed to crack bridging/plastic stretching of the ductile phase [15, 39, 40], and crack deflection mechanisms [13]. Soboyejo et al. [2] concluded that toughening in this system is primarily due to continuous tilting and twisting of the crack paths around the niobium particles. Fracture in this system occurs by the transgranular cleavage mode [2] at room temperature. Toughness variation as a function of relative density for the matrix reinforced with coarse Nb powder is shown in Fig. 2. MoSt2 reinforced with 20 vol.% Nb particles (75-180 ~m) showed the maximum resistance to crack propagation among all the materials in this study. At low densities (93-95%) KQ remained close to 5 MPax/-m and increased to ~ 5.7 M P a x / m at higher densities. This is approximately a 35-40% increase in KQ compared to the unreinforced matrix. Fracture surface morphologies of the MoSiz-Nb composite consolidated at 1200°C for 1 h [Fig. 3(c)] and the MoSi2-Nb sample consolidated at 1400°C for 1 h [Fig. 3(d)] show that matrix fracture is combined transgranular and intergranular in nature. The Nb particle undergoes brittle cleavage which means that crack blunting and crack bridging mechanisms are unlikely in this material. This is shown in Fig. 3(c). The brittle cleavage fracture of ductile Nb is an unlikely but interesting feature and has been observed previously [29]. Experiments on a lead-glass system by Ashby et al. [44] have shown that the yield stress of a constrained material is higher than without any constraints. In a similar manner, the matrix may be viewed as constraining the plastic zone near the crack tip, causing a plane strain state, leading to brittle failure. The MoSi2-Nb reaction zone may cause the initial crack length in the Nb particle to be high enough, resulting again in a plane strain state. It leads

SURYANARAYANAN et al.: MECHANICAL PROPERTIES OF MoSi:

3755

Fig. 3. Fracture surface morphologies of monolithic and composite MoSi2. Consolidations conditions were: (a) MoSi2--1200°C, l h, (b) MoSi2--1400°C, 4h, (c) MoSi2-Nb (coarse~120@'C, l h; (d) MoSi2 Nb (coarse)--1400°C, 1 h; (e) MoSi2-SiC 120@'C, 1 h, and (f) MoSi2-SiC--1400°C, 4 h. Pressure in all cases was 207 MPa. Arrows point to cusp shaped pores.

to the conclusion that optimal HIP processing alone will not be sufficient for the requisite toughness improvement of this composite. In addition to optireal HIPing, the MoSi2/Nb interface will need to be engineered to enhance its effectiveness in toughening, for example, by coatings, It must be realized that there is an inherent limitation to the toughening increase that is possible with spherical reinforcements. This is because an advancing crack has a high probability of entirely avoiding

the reinforcements, due to the relatively small deviation required to circumvent the spheres [29]. 3.3.3. MoSie-2Ovol.% SiC. Figure 2 shows the toughness increase in MoSi2-SiC with increasing relative density. The resistance to crack propagation in this case is considerably lower than for monolithic and Nb reinforced MoSi2. Previous studies of fracture toughness on MoSi2 reinforced with SiC whiskers have produced materials with higher fracture toughness values than the base matrix

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SURYANARAYANAN et al.: MECHANICAL PROPERTIES OF MoSi2

[8, 13, 14, 45, 46]. Carter and Hurley [45] reported a 54% increase in fracture toughness in this material compared with the unreinforced matrix. It was also found that cracks are generally intergranular. The mechanisms which contribute to toughening in this material are crack deflection, crack bridging, microcracking, and crack bowing [45]. The fracture toughness values in the present study are considerably lower than in earlier studies. The primary reason for this difference is that the samples in the present study have been consolidated to much lower levels of density than those in previous studies where samples were consolidated to near theoretical densities. As mentioned in Ref. [28] SiC whiskers actually inhibit densification and therefore higher HIP temperatures or pressures are required to achieve complete densification, Fracture surface analysis on MoSi2-SiC samples consolidated at 1200°C [Fig. 3(e)] shows intergranular fracture in the matrix. The density in this material is very low, resulting in easy puilout/decohesion of the SiC whiskers. Fracture topologies for the sample HIPed at 1400°C for 4 h is shown in Fig. 3(f). Figure 3 (f) also shows the cleavage mode of fracture of the matrix and the brittle fracture in the SiC whiskers. An interesting microstructural feature in the SiC-reinforced material is that there is a clumping of SiC whiskers that are either rod shaped or stacked equiaxed SiC grains [28]. As mentioned earlier, this agglomeration produces highly brittle regions (in comparison to the matrix), which may be another reason for the low fracture resistance of the SiC reinforced materials in this study. If electro-static attractions are the reason for agglomeration, a feasible method to remove these attractive forces can be to coat the whiskers by a suitable material. If the material is chosen such that it is ductile and has low reactivity with the matrix and SiC, it will also act as a ductile reinforcement enhancing the fracture resistance of the composite. Agglomeration can be eliminated by ultrasonically mixing constituents in liquid medium or multi-stage milling processes [47]. The intergranular mode of fracture in the low density sample may be due to the fact that matrix porosity level is high enough to make the grain boundaries much weaker than the crystal cleavage planes. The general behavior of KQ vs relative density in this case is similar to the other two cases investigated. The one surprising distinction is that the 1 h samples have a higher KQ than the 4 h samples. This is due to the bigger agglomerate sizes and, therefore, a larger spacing between agglomerate regions in t h e specimens consolidated for 4 h [28]. Also, there is a higher probability of the agglomerate region being at the crack tip. These factors explain the lower resistance to crack propagation in the samples consolidated for 4 h.

The sudden increase in toughness values after reaching density values greater than 90% is similar to the conclusions of Fleck and Smith [48]. It must be

emphasized that none of these tests are valid planestrain fracture toughness tests and the results, although qualitatively correct, cannot be used in actual design calculations. Apart from the fact that fatigue precracking was not done, the number of tests were not adequate to statistically neutralize the effects of surface roughness, non-standard varying sizes of test specimens, and the possibility of a niobium or silicon carbide particle at the crack tip. Some or all of these reasons may lead to different levels of error in these results. 4. CONCLUSIONS 1. Yield strengths of monolithic MoSi 2 were found to be greater than that for MoSi2-Nb, which were in turn greater than those for MoSi2-SiC, at test ternperatures of 1100 and 1300°C. The low strength of the SiC reinforced material is probably due to the higher level of porosity in the material. The larger level of grain boundary silica phase in this composite, which may be viscous at high temperatures, is probably another reason for the low strength measured for this material. 2. At a particular applied stress, the creep rates are higher in MoSi2-Nb than in monolithic MoSi2, and MoSi2-SiC. The PLC stress exponents for monolithic MoSi2 and MoSi2-SiC are comparable to the values obtained in other studies. The PLC stress exponent for MoSi2-Nb is 2.2, with a PLC activation energy of 160kJ/mol. 3. Room temperature fracture is predominantly cleavage for MoSi2 and MoSi2-Nb samples with densities close to theoretical values. Toughness values are almost constant until a relative density of ~ 9 5 % is achieved. This level of density corresponds closely to the range where there is a change in the nature of porosity from being interconnected (cusp shape) to being isolated (spherical shape). However, in all three cases, as relative density changes from ~95 to 100%, there is a disproportionate increase in toughness values. 4. MoSi2-Nb was found to have the maximum resistance to crack propagation, followed by MoSi 2 and MoSi2-SiC, in that order. The uncharacteristically low fracture toughness of MoSi2-SiC is probably due to the relatively low density level (primarily interconnected porosity) in this case compared to the other two and/or because of the agglomeration of the SiC whiskers. Acknowledgements--This research was conducted under an

AFOSR contract with Dr A. H. Rosenstien as program manager. The help of Mr Bernie Sunier with mechanical testing is acknowledged. REFERENCES 1. A. K. Vasudevan and J. J. Petrovic, High Temperature Structural Silicides. Elsevier Science, Amsterdam, (1992).

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