Intermetallics 14 (2006) 1461e1471www.elsevier.com/locate/intermet

Microstructure and mechanical behaviour of reaction hot pressedmultiphase MoeSieB and MoeSieBeAl intermetallic alloys

R. Mitra a,*, A.K. Srivastava a,b, N. Eswara Prasad b, Sweety Kumari b

a Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, Indiab Defence Metallurgical Research Laboratory, Hyderabad 500 058, India

Received 22 September 2005; received in revised form 9 January 2006; accepted 30 January 2006

Available online 4 April 2006

Abstract

Microstructures of 76Moe14Sie10B, 77Moe12Sie8Be3Al, and 73.4Moe11.2Sie8.1Be7.3Al alloys, processed by reaction hot pressingof elemental powder mixtures, have shown a-Mo, Mo3Si, and Mo5SiB2 phases. In addition, particles of SiO2 formed from the oxygen content ofraw materials could be seen in the 76Moe14Sie10B alloy, while a-Al2O3 formed in the alloys containing Al. Parts of the Al have been foundwithin the solid solutions of a-Mo and Mo3Si. The average fracture toughness determined from indentation crack lengths and three-point bendtesting of single edge notch bend specimens lies in the range of 5.0e8.7 MPaOm, with alloys containing Al demonstrating higher values. Anal-yses of load-displacement plots, fracture profiles and indentation crack paths have shown evidence of R-curve type behaviour and operatingtoughening mechanisms involving crack bridging by a-Mo, crack deflection and branching. Flexural strength is related to volume fraction ofthe a-Mo and Al content. Compression tests on the 76Moe14Sie10B alloy between 1100 �C and 1350 �C have shown excellent strength re-tention, and evidence of thermally activated plastic flow.� 2006 Elsevier Ltd. All rights reserved.

Keywords: A. Molybdenum silicides; A. Multiphase intermetallics; B. Fracture toughness; B. Mechanical properties at high temperature; C. Reaction synthesis

1. Introduction

The quest for materials to be applied in aero-engine struc-tural components operating in the temperature range higherthan that of Ni-base superalloys, that is 1100e1500 �C inair or oxidizing environments, has triggered extensive interestin the research and development of molybdenum silicides. TheMoeSi binary phase diagram [1] contains three intermetalliccompounds, MoSi2, Mo5Si3 and Mo3Si with melting pointsof 2030 �C, 2180 �C and 2025 �C, respectively. MoSi2 isknown for its outstanding oxidation resistance, because ofthe formation of a protective and impervious film of SiO2

[2,3]. On the other hand, since the SiO2 scale on Mo5Si3 orMo3Si is not impervious and protective, those are prone to

* Corresponding author. Tel.: þ91 3222 283292; fax: þ91 3222 28280.

E-mail address: rahul@metal.iitkgp.ernet.in (R. Mitra).

0966-9795/$ - see front matter � 2006 Elsevier Ltd. All rights reserved.

doi:10.1016/j.intermet.2006.01.057

damage by oxidation at high temperatures [4]. It is also wellestablished that the room temperature ductility of polycrystal-line molybdenum silicide based intermetallic alloys and com-posites is negligible, and the fracture toughness is poor [5,6].Composites with brittle or ductile reinforcements have shownonly modest rise in fracture toughness [7,8].

In the late 1990s, it was reported that alloying of Mo5Si3with boron significantly improves the oxidation resistance,due to the formation of a borosilicate glassy film possessinghigher fluidity compared to that of SiO2 and its self-healingcharacter [9]. In addition, the creep strength is much higherthan that of MoSi2 [4,10]. Subsequently, interest has grownin the development of multiphase MoeSieB alloys [11e18],with a composite microstructure comprising a-Mo, Mo3Siand Mo5SiB2 (also known as T2) phases. Fig. 1 shows theMo-rich section of the ternary, isothermal MoeSieB phasediagram for 1600 �C, originally contributed by Nowotnyet al. [19]. a-Mo has a bcc structure with the solubility for

1462 R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

Si atoms up to 3 at.% and less than 1 at.% solubility forB atoms, while Mo3Si possesses A15 (cubic) structure withsingle-phase composition close to 76Moe24Si (at.%) [20],suggesting deviation from stoichiometry. The room tempera-ture fracture toughness of Mo3Si has been reported to beclose to 3.0 MPaOm. On the other hand, Mo5SiB2 possessesa body-centred tetragonal structure with 32 atoms in theunit cell, comprising 20 atoms of Mo, 4 atoms of Si and 8atoms of B. Mo5SiB2 single crystals have shown extremelybrittle behaviour with a fracture toughness of about 2.0MPaOm, which is less than that reported for MoSi2 basedmaterials [5e8]. However, the creep strength of Mo5SiB2 ismuch superior to that of MoSi2 [21]. Presence of coarseand interconnected a-Mo particles in the MoeSieB alloyshas been reported to lead to a significant improvement infracture toughness by crack bridging mechanism [18]. A thor-ough review of the mechanical behaviour of silicides, includ-ing those of the MoeSieB alloys may be found in Ref. [22].Only a limited understanding exists on the structureepropertyrelations and high temperature deformation behaviour of theMoeSieB alloys, particularly those made by powder metal-lurgy processing.

In the present investigation, microstructure and mechanicalproperties of a ternary MoeSieB alloy and two quaternaryMoeSieBeAl alloys, processed by reaction hot pressing ofelemental powders have been studied. The emphasis is on un-derstanding the fracture mechanisms and high temperature de-formation behaviour of the selected alloys. Al has been chosenas alloying element, so that the oxygen present in the powdermixture is scavenged to form dispersoids of Al2O3 in situ dur-ing processing. In previous investigations [23,24], Al additionto the powders of MoSi2 or mixtures of Mo and Si powders,which were subsequently hot pressed, has led to the formationof a-Al2O3 by possible in situ reactions as follows:

Fig. 1. Mo-rich section of the MoeSieB ternary phase diagram [19], with ‘‘*’’

showing the location of the 76Moe14Sie10B alloy.

3SiO2ðsÞ þ 4Alðs or lÞ/ 2Al2O3ðsÞ þ 3SiðsÞ ðiÞ

MoO3ðsÞ þ 2Alðs or lÞ/ Al2O3ðsÞ þ MoðsÞ ðiiÞ

B2O3ðsÞ þ 2Alðs or lÞ/ Al2O3ðsÞ þ 2BðsÞ ðiiiÞ

The reactions (i)e(iii) are driven by reduction in the Gibbsfree energy [25], because the free energy of formation ofAl2O3 is lower, compared to those of SiO2 and MoO3.

2. Experimental procedure

The raw materials used in the present study are Mo, Si, Band Al powders. The source and level of purity of Mo, Si, Band Al powders are shown in Table 1. The powder mixturesof elemental Mo, Si, B, and Al were weighed, so as to obtainthe chemical compositions of 76Moe14Sie10B (MSB),77Moe12Sie8Be3Al (MSB3Al), and 73.4Moe11.2Sie8.1Be7.3Al (MSB7.3Al), and dry blended for 24 h insideagate pots in a planetary ball mill. Subsequently, the mixedpowders were wrapped in grafoil, and compacted in argon en-vironment inside graphite dies with a pressure of 26 MPa at1500 �C for 1 h, and annealed at 1600 �C for 0.5 h. Prior tohot pressing, the MoeSieBeAl powder mixtures were heldat 700 �C for 1 h in order to allow the liquid Al to reactwith SiO2, so that the particles of a-Al2O3 could form insitu. A higher temperature would have increased the fluidityof liquid Al, and allowed it to ooze out of the die assembly.The hot pressed products were discs with a diameter of75 mm and thickness of 4e5 mm. Graphite was removedfrom the top layers through abrasive action of sand blasting,followed by polishing on coarse and fine diamond-coateddiscs. The hot pressed discs were sectioned using a dia-mond-coated saw or wire-cut electro-discharge machining(EDM) for examination of the microstructure and conductingmechanical tests. The phases were identified using X-Ray dif-fraction (XRD). The microstructures were observed on thescanning electron microscope (SEM) and the electron probemicro-analyzer (EPMA) using both secondary (SE) and back-scattered electron (BSE) imaging modes. The volume fractionsof a-Mo and Mo3Si were determined using point countingtechnique on image analyzer. Energy Dispersive X-ray

Table 1

Source, purity levels and particle size of raw materials used

Raw

material

Source Particle

size (mm)

Purity

level (%)

Impurities (wt.%)

Mo Plansee, Austria 4.4e5.2 99.5 O z 0.1, Fe, C, Ni, Cr,

Al, Cu, K, W, Ca, Si,

Mg, Na, Ti� 0.01

Si Johnson Matthey,

USA

20 99.8 Fe z 0.07, C, Mg, O,

N� 0.01

Al NALCO,

Bhubaneswar,

India

40 99.5 O z 0.29, Fe z 0.05,

Si z 0.05, Cu, Zn,

Mg� 0.002

B Advanced Research

Centre, Hyderabad,

India

5 98 Si z 0.18, C z 0.772,

O z 1 wt.%, N� 0.01.

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Analysis (EDAX) on SEM or Wavelength Dispersive Spec-troscopy (WDS) on EPMA were used to estimate the chemicalcompositions of various constituent phases.

Microhardness of the individual phases was determinedusing a Vickers diamond indenter, operated with a loadof 0.098 N. An average was calculated from the results of aminumum of five microhardness tests. The micro-indentations,made on relatively coarse a-Mo particles, had average diagonallengths in the range of 4.6e4.8 mm, such that the cornerslay well within the phase-limits. The lengths of diagonals ofindentation inside the Mo3Si phase ranged between 3.3 and3.6 mm, while those inside the Mo5SiB2 phase varied between3.1 and 3.3 mm. The bulk Vickers hardness was also deter-mined using loads between 98.1 N and 294.3 N. Flexural testswere conducted on specimens, having a cross-section of 4 mm(¼ b)� 3 mm (¼ d ), were tested on INSTRON 8801 ata crosshead displacement rate of 0.5 mm per min usingthree-point bend fixtures of SiC with a span (S ) of 40 mm.The flexural strength, sB has been determined from the resultsof three-point bend tests using the relationship [26]:

sB ¼ 3PmaxS=2bd2 ð1Þ

where Pmax is the maximum load recorded prior to failure.Fracture toughness tests were conducted using the same

set-up as that for flexural tests on single edge notch bend(SENB) specimens with 4 mm thickness and 8 mm width,and having a notch of 4 mm depth machined with the helpof wire-cut EDM. The fracture toughness, KQ has been deter-mined using the relationship [27]:

KQ ¼�PmaxS=BW3=2

�f ða=WÞ ð2Þ

where B is the thickness, W is the width, and a is the depthof the notch in the direction of width. For three-point bendtesting, f(a/W )¼ [3(a/W )1/2[1.99� (a/W ){1� (a/W )}{2.15� 3.93(a/W )þ 2.7(a/W )2}] / [2{1þ (2a/W )} {1� (a/W )}3/2].Three specimens of each of the alloys were tested for measure-ment of the average flexural strength and SENB fracturetoughness. In addition, the crack lengths at the corners ofVickers hardness indentations were measured, and the fracturetoughness was calculated using the empirical relationship pro-posed by Evans and Charles [28]:

KI ¼ 0:16Hffiffiffiapðc=aÞ�3=2 ð3Þ

where H is the hardness, c is the characteristic crack lengthand a is one-half of the diagonal length. The characteristiccrack length was determined from length, l of the cracks em-anating from the indentation corner to its end point, using therelation:

c¼ lþ a ð3aÞ

Evans and Charles have used stress analysis and curve fittingtechnique to derive the relationship for determination of theindentation fracture toughness. Eq. (3) is valid for mediancracks, in which the c/a ratio is equal to 2.5 or more. In

the present study, cracks for loads of 30 kg or more werefound to be median in character, and were used for calcula-tion of fracture toughness. One of the advantages of usingEq. (3) is that the value of Young’s modulus, which is un-available for the alloys studied, is not required. Young’smodulus data are required for some of the other empiricalequations [29e32].

Compression tests were conducted on the EDM wire-cutcylindrical specimens of the MSB alloy, having dimensionsof 2.8 mm diameter and 4 mm height at nominal strain ratesof 10�3 s�1 or 10�4 s�1 between 1100 �C and 1350 �C inair. Damage due to oxidation was minor and did not seem toaffect the tests. Similar tests could not be performed on the al-loys containing Al, because of severe oxidation accompaniedby mass loss during heating, which would have affected theproperties, even if a stable and protective oxide scale is ex-pected to form after a while.

3. Results and discussion

3.1. Microstructure

The composition and phase field of the MSB alloy isshown with asterisk in Fig. 1 [19]. Fig. 2(a) and (b) showsan SEM BSE image of the MSB alloy at lower magnificationand an EPMA BSE image at higher magnification, respec-tively. The phases present in the microstructure, a-Mo,Mo3Si and Mo5SiB2, exhibit differential contrast in theBSE images [Fig. 2(b)], recorded using the EPMA. The pres-ence of a-Mo, Mo3Si and Mo5SiB2 phases has also beenconfirmed from the peaks of XRD pattern [Fig. 2(c)], whichis consistent with the expectations from the phase diagram(Fig. 1). The irregularly shaped a-Mo particles or clustersof those have sizes ranging between 5 and 15 mm. ThroughWDS analysis, the a-Mo has been found to contain 2 at.%Si in solid solution. The concentration of B in solid solutiondetermined indirectly, from the difference between unity andsum of atom fractions of Mo and Si, to be 1 at.% is at bestapproximate, and is close to that mentioned by Liu et al.[13]. Error also arises due to the possibility of presence ofsmall concentration of O in solid solution, originating inMo powders used as raw material (Table 1). In the ternaryphase diagram for 1600 �C reported by Nunes et al. [33],a-Mo has been shown as containing about 1 at.% B and2e3 at.% Si. However, more recently, ternary phase diagramof Sakidja and Perepezko has shown negligible solubility ofB in Mo [34,35]. On the other hand, Schneibel [36] hasreported the concentration of B as 3.2 at.% in the a-Mophase present inside the arc-melted ingots of Moe19.5We12Sie8.5B alloy, which could be due to the effect of W insolid solution with Mo. Probably further experiments are re-quired to resolve the issue of solubility of B in a-Mo. Thedistribution of a-Mo (appearing white) in Fig. 2(b) appearsclose to uniform at lower magnification in Fig. 2(a), andits volume fraction has been found to be 33%. The oxygenand nitrogen content in the MSB alloy has been found usingthe vacuum fusion technique to be 0.44 wt.% and

1464 R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

Fig. 2. MSB alloy: BSE micrographs recorded on (a) SEM at lower magnification, and (b) EPMA at higher magnification, showing the location of a-Mo, Mo3Si

and Mo5SiB2 phases; and (c) XRD pattern.

0.006 wt.%, respectively. The oxygen in the hot pressedproduct is attributed to that contained in the powders(Table 1), and is present in the form of silica (SiO2) particles.SiO2 also forms, because the Si reacts with the oxygen con-tained in the powders of Mo and B. Fig. 3 is the EPMA BSEimage showing SiO2 particles, the presence of Si and O inwhich has been confirmed through WDS X-ray elementalmapping. SiO2 has been found to account for about 9vol.%. The actual amount of Si involved in forming Mo3Siand Mo5SiB2 during reaction hot pressing is reduced tosome extent, because of the formation of SiO2.

Fig. 4(a) presents the EPMA BSE image, showing themicrostructure of MSB3Al alloy, while Fig. 4(b) is the WDSX-Ray elemental map of Al, showing the distribution of a-Al2O3

particles in the same microstructure, accounting for about 10vol.%. In the alloys containing Al, the phase composition ofmicrostructure is expected to differ from that in the MSB al-loy, because (i) a significant fraction of Al reacts with SiO2

to form a-Al2O3 and release free Si, and (ii) only a part ofthe Al enters a-Mo and Mo3Si to form solid solution. Assum-ing the oxygen content to be same as that of the MSB alloy,and that the entire SiO2 is converted to Al2O3, simple calcula-tions using appropriate chemical equations show that 1.6 at.%and 6.2 at.% Al must be available for alloying in the MSB3Aland MSB7.3Al alloys, respectively. Again, the released Si isexpected to react with the Mo, and increase the volume frac-tion of Mo3Si at the expense of a-Mo.

Fig. 5(a) and (b) show the EPMA BSE images of the micro-structure of the MSB7.3Al alloy at two different magnifica-tions. The micrograph at lower magnification shows thatthe distribution of a-Mo (appearing bright) is non-uniform,while Fig. 5(b) represents a region with high volume fraction.The distribution and volume fraction of the a-Al2O3 in theMSB7.3Al alloy, examined through WDS X-ray mapping,have been found to be similar to that in the MSB3Al alloy.The characteristic XRD pattern from the MSB7.3Al alloy is

Fig. 3. EPMA BSE image of MSB alloy showing SiO2 particles dispersed.

1465R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

Fig. 4. MSB3Al alloy: (a) EPMA BSE image; and (b) WDS X-ray map of Al.

presented in Fig. 5(c), which shows the peaks of a-Mo, Mo3Si,Mo5SiB2, and a-Al2O3. It may be noted that the intensity orheight of the Mo (110) peak is more than that of the Mo3Si(210) peak in the XRD pattern of the MSB alloy sample[Fig. 2(c)], while it is opposite in the case of the MSB7.3Alalloy [Fig. 5(c)]. The trend observed for the XRD peak inten-sities is consistent with the lower volume fraction of a-Mo inthe MSB7.3Al alloy, in which the atomic concentration of Mohappens to be the lowest. The volume fractions of a-Mo havebeen determined to be 35% and 29% in the MSB3Al andMSB7.3Al alloys, respectively. Part of the Al goes into solidsolution into the a-Mo and Mo3(Si1�xAlx) phases of the

MSB3Al and MSB7.3Al alloys. While the concentration ofAl in a-Mo has been determined in the range of 1e2 at.%,the value of x in the Mo3(Si1�xAlx) phase has been found tobe 0.24 and 0.36 in the MSB3Al and MSB7.3Al, respectively.B has not been detected in the Mo3(Si,Al). Isolated inclusionsof coarse Mo-boride particles [appearing dark in Fig. 5(a)]could be found in the MSB7.3Al alloy.

3.2. Room temperature mechanical properties

Table 2 shows the bulk hardness, flexural strength and frac-ture toughness of the MSB, MSB3Al and MSB7.3Al alloys,

Fig. 5. MSB7.3Al alloy: EPMA BSE image at (a) low and (b) high magnifications, and (c) typical XRD pattern from the MSB7.3Al alloy, showing the peaks of

a-Mo, Mo3(Si,Al), Mo5SiB2 and a-Al2O3.

1466 R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

along with the average microhardness of a-Mo in each. Themechanical properties of the reaction hot pressed (RHP)MoSi2 [37] with average grain size of 5 mm have been shownin Table 2 for comparison.

3.2.1. HardnessThe microhardness of a-Mo, Mo3Si, and Mo5SiB2 phases

in the MSB alloy has been found to be equal to or in the rangeof 8.4 GPa, 14e16 GPa, and 18e20 GPa, respectively, andagree well with the values cited in literature [38]. The valuesof microhardness of a-Mo in the MSB and MSB3Al alloysare similar, while it is slightly lower in the MSB7.3Al alloy.The measured microhardness of a-Mo is significantly higherthan that of bulk molybdenum in annealed or wrought form(1.7e2.6 GPa depending on whether it is annealed or rolled)[39], which could be due to one or more of the following,namely, the constrained plastic deformation involving a triaxialstate of stress, marginal interference of underlying or sur-rounding hard intermetallic phases, Mo3Si and Mo5SiB2, or in-terstitial oxygen or boron content. The bulk hardness of thealloys does not appear to vary very significantly with compo-sition only, since alterations in volume fractions of phases alsoneed to be taken into account. Hardness of the MSB3Al alloyis lower than that of the MSB and MSB7.3Al alloys by closeto 10% on account of its higher volume fraction of a-Mo.

3.2.2. Flexural strengthThe flexural strengths of the Al containing alloys are less

than those of the MSB alloy, with MSB7.3Al showing signif-icantly lower value. Schneibel et al. [15] have reported flexuralstrengths in the range of 380e410 MPa for the Mo12Si8.5Balloy, which is similar to those obtained for the MSB andMSB3Al alloys in the present study. The flexural strengthsof MSB and MSB3Al alloys are about twice that of MoSi2,probably due to the contribution of much harder phases,Mo3Si and Mo5SiB2, as well as the constrained plastic defor-mation of a-Mo. While solid solution of Al could have par-tially contributed to the low flexural strength of MSB7.3Alalloy, non-uniform distribution and lower volume fraction ofthe ductile a-Mo phase compared to that of brittle phases isalso responsible. The coarse and brittle inclusion of Mo-borides in the MSB7.3Al alloy could also have reduced theflexural strength.

Table 2

Mechanical properties of alloys A, B, and C and their comparison with data on

reaction hot pressed MoSi2 [37]

Material Hardness

(GPa)

Microhardness

of a-Mo

(GPa)

Flexural

strength

(MPa)

Fracture toughness

(MPaOm)

SENB Indentation

MSB 9.9� 0.5 8.4� 0.3 404� 10 5.0� 0.8 5.2� 1.0

MSB3Al 8.7� 0.2 8.4� 0.2 337� 7 5.8� 1.4 7.2� 1.5

MSB7.3Al 10.0� 0.6 8.4� 0.2 167� 15 6.6� 1.5 8.7� 1.7

RHP MoSi2 9.3� 0.1 N.A. 193� 10 4.8� 0.1 5.0� 0.3

‘‘SENB’’ refers to fracture toughness determined through testing of Single

Edge Notch Bend specimens. N.A. stands for ‘‘Not Applicable’’.

3.2.3. Fracture toughness3.2.3.1. SENB tests. The fracture toughness obtained fromtesting of the SENB specimens has been found to be in therange of 5.0e7.5 MPaOm, with the MSB3Al and MSB7.3Alalloys showing an average improvement of about 15 and30%, respectively, with respect to that recorded for the MSBalloy (Table 2). The scatter in the results of fracture toughnesstests is probably because of the heterogeneous character of mi-crostructure, and is similar to the observations previously re-ported for Mo15.3Si11.5B alloy [40]. Choe et al. [17,38]have reported fracture toughness of 4.1 MPaOm for P/M pro-cessed Mo16.8Si8.4B alloy, and 7.2 MPaOm for arc-meltedMo12Si8.5B alloy. Hence it may be inferred that the data ob-tained through fracture toughness testing in the present studybroadly fall in the range reported in the previous publications[15,17,38]. It is obvious that the fracture toughness of theMSB, MSB3Al and MSB7.3Al alloys are marginally or signif-icantly higher compared to that of single-phase MoSi2 (z4e5 MPaOm) [7], Mo5Si3 (2e3 MPaOm) [41,42], Mo3Si(z3 MPaOm) [19], and Mo5SiB2 (z2 MPaOm) [43]. Com-pared to the fracture toughness of 3.3 MPaOm shown by anMo3Sie50 vol.% T2 composite (no a-Mo) [44], the data re-corded in the alloys of the present study are certainly an im-provement, emphasizing the role of somewhat ductile a-Modispersion.

Fig. 6 shows a typical load-displacement plot of theMSB3Al alloy, demonstrating evidence of deviation from lin-ear elastic fracture, involving typically non-catastrophic crackpropagation. The crack propagation characterized initially bya drastic fall in load after its peak is reached, is interruptedby increase in load with increasing displacement, which isagain followed by a gradual decrease in load. In other words,the process of fracture is relatively stable and graceful, drivenby the energy supplied externally. However, not all the load-displacement plots for SENB specimens of the MSB3Al andMSB7.3Al alloys have shown such deviation from linear elas-tic behaviour, suggesting that position of notch in the micro-structure and chance encounter with the phases obstructingcrack growth are the essential criteria for graceful fracture.Study of the fracture profiles by placing the broken halvesof the SENB specimens of the MSB3Al alloy together inside

0.000.00 0.05 0.10 0.15

0.05

0.10

0.15

Lo

ad

kN

Displacement mm

Fig. 6. Typical plot of load against displacement from fracture toughness test

conducted on the SENB specimen of MSB3Al alloy, showing evidence of in-

terruption of crack propagation and graceful failure.

1467R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

the SEM has shown distinct evidence of deflection and branch-ing of cracks [Fig. 7(a) and (b)] in the course of crack propa-gation. The non-catastrophic failure involves crack deflection,branching [Fig. 7(a) and (b)] and plastic deformation alongwith strain hardening of ductile a-Mo phase. The top view im-ages of the fracture surfaces show a-Mo particles (confirmed byEDAX analyses), as well as cleavage facets [Fig. 8(aed)].Evidence of mixed-mode failure involving shear has alsobeen noticed in the fracture surfaces [Fig. 8(b) and (c)], partic-ularly for regions having a higher volume fraction of a-Mo.Mixed-mode failure is expected because of the differences inthe Young’s modulus and yield strength or hardness of theconstituent phases, a-Mo, Mo3Si and Mo5SiB2. As a result,a crack would experience sharp changes in stress intensity fac-tor on passing from one phase to another, and will tend to fol-low the easiest path.

3.2.3.2. Indentation tests. Table 2 shows that the indentationfracture toughness of the MSB alloy is close to that deter-mined by testing of SENB specimens. On the other hand, in-dentation fracture toughness of MSB3Al and MSB7.3Alalloys is close to the upper limit of the scatter band of data ob-tained from testing of the SENB specimens. The larger differ-ence in fracture toughness obtained through indentation andtesting of SENB specimens in the case of the MSB3Al orMSB7.3Al specimens, compared to that in the case of theMSB alloy, may be explained because of the higher degreeof non-uniformity in distribution of a-Mo in the Al-alloyedsamples, and operation of toughening mechanisms affectingthe short cracks at corners of indentations.

Studies of the indentation crack paths have shown clearevidence of crack arrest and bridging by the relatively ductile

a-Mo particles in the MSB alloy [Fig. 9(a)]. Deflection andbranching of indentation cracks have been observed exten-sively in the MSB3Al and MSB7.3Al alloy specimens, asshown in Fig. 9(b) and (c), respectively. The indentation crack-ing behaviour may be analyzed further, assuming that the char-acteristic crack length, c is related to load, P by a power-lawrelation:

c¼ kPm ð4Þ

where k is a constant, and m is an exponent. The m can befound from the slope of the best-fit line obtained through thelinear regressional analysis of the data of log c against log P,since,

log c¼ log kþm log P ð4aÞ

Eq. (4a) leads to:

dðlog cÞ=dðlog PÞ ¼ m

Thus the rate of growth of crack with increasing load, dc/dP isgiven by:

dc=dP¼ mðc=PÞ ð4bÞFig. 10 displays the plots depicting the variation of loga-

rithm of crack length against the logarithm of load. The cracklengths of the MSB3Al and MSB7.3Al alloys are shorter, com-pared to those of the MSB alloys, and are consistent with thetrends observed in fracture toughness testing. The values of mhave been found to be 0.89 for MSB alloy, 0.56 for MSB3Alalloy, and 0.50 for MSB7.3Al alloy. The observation, that thevalues of m for all the alloys are less than 1.0, suggests that the

Fig. 7. SEM images of the profile of the fracture surface of the MSB3Al alloy, presenting the evidence of: (a) crack deflection (arrowed) and (b) branching.

1468 R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

Fig. 8. Typical SEM images of the fracture surfaces of: (a) and (b) MSB alloy, (c) MSB3Al alloy, and (d) MSB7.3Al alloy. The fracture surfaces show a-Mo

particles appearing bright.

rate of increase of crack length with load is less than that inperfectly brittle materials. A linear relationship between cracklength and applied load has been reported for basically brittleglass ceramics [45]. It is also clear that the values of m for the

MSB3Al and MSB7.3Al alloys are much less than that ofthe MSB alloy. Smaller m implies that a higher load isrequired for generation of cracks of similar length in theMSB3Al or MSB7.3Al alloys.

Fig. 9. Paths of the cracks originating at the corners of indentations and undergoing (a) arrest or bridging at a-Mo particles (arrowed) in the MSB alloy, deflection

or bridging in the (b) MSB3Al and (c) MSB7.3Al alloys.

1469R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

The rates of change in crack lengths with load, dc/dP havebeen plotted with respect to the applied load in Fig. 11. Thedecrease in dc/dP with increasing P is clear and significantfor the MSB3Al or MSB7.3Al alloys, and only marginal andless obvious in the case of the MSB alloy. The value of dc/dP for a specific load is the highest for the MSB alloy andthe least for MSB3Al alloy. In other words, the indentationcracking behaviour demonstrates an R-curve type behaviour,in which the resistance to fracture increases with increasingcrack length. With increase in the length of a crack, a greaternumber of microstructural obstacles are encountered, which isresponsible for the reduced rate of crack growth with increas-ing load.

The shorter indentation cracks in the case of the MSB3Aland MSB7.3Al alloys, compared to those recorded for theMSB alloy at a particular load, is in tune with the results of

-4.11.9 2.0 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8

-4.0

-3.9

-3.8

-3.7

-3.6

-3.5

Lo

g (C

har

acte

rist

ic C

rack

Len

gth

, m)

Log (Load, N)

MSBMSB3ALMSB7.3AL

Fig. 10. Plots of logarithm of indentation crack lengths against the logarithm

of load for the MSB, MSB3Al and the MSB7.3Al alloys.

100 200 300 400 5002

3

4

5

6

7

8

9

dc/

dP

(X

107

m/N

)

Indentation Load (N)

MSBMSB3ALMSB7.3AL

Fig. 11. Plots showing the rate of change of indentation crack length with re-

spect to load for MSB, MSB3Al and MSB7.3Al alloys.

tests on SENB specimens (Table 2). The toughening mecha-nism in the MSB alloy is primarily due to the bridging and ar-rest of the indentation cracks by the ductile a-Mo particles[Fig. 9(a)]. On the other hand, the profiles of crack paths inthe MSB3Al and MSB7.3Al samples after fracture toughnesstests [Fig. 7(a) and (b)] and indentation cracking [Fig. 9(aec)]show sharp deflections by large angles (60 � or more), in addi-tion to bridging and branching, leading to greater tortuosityand toughening. The Al2O3 particles are expected to promotecrack deflection as well.

It may be postulated that the fracture toughness of theMSB3Al and MSB7.3Al alloys is higher than that of theMSB alloy, partly because Mo3(Si,Al) may not be as brittleas Mo3Si on account of reduction in the covalent characteror directionality of bonding, and increase in the metallic com-ponent of the bond, as has been proposed for MoSi2 throughtheoretical analyses by Waghmare et al. [46]. However, furtherexperimental work is required to be carried out on single-phase Mo3(Si,Al) to confirm that since the present situationis more complicated due to the multiphase microstructureand flaws typical of PM products.

3.3. High temperature compression tests

The yield stress and maximum stress of the MSB alloys ob-tained from the high temperature compression tests in therange of 1100e1350 �C are shown in Table 3. For a fixed tem-perature, the flow stress is higher on testing at the strain rate of10�3 s�1, when compared to that corresponding to the strainrate of 10�4 s�1, which is normally anticipated. The maximumstress, the corresponding strain for which defines the limit ofuniform deformation, too varies with temperature and strainrate in the same way as the yield stress does. The data asa whole suggest impressive strength retention due to the pres-ence of Mo5SiB2 and Mo3Si phases. Significant deformationof Mo3Si and Mo5SiB2 is ruled out, since the brittle-to-ductiletransition temperatures are at 1400 �C or 1500 �C, respec-tively [20,40]. It is clear that the yield stress decreases from950 MPa at 1100 �C to around 280 MPa at 1350 �C, whentested at the strain rate of 10�4 s�1. The decrease in yieldstress with increasing temperature is significant between1100 �C (1373 K) and 1200 �C (1473 K), which may be ex-plained based on the fact that 1373 K and 1473 K correspondto 0.48Tm and 0.51Tm, respectively where Tm is the absolutemelting point (2890 K) of Mo. Again, the a-Mo is expected

Table 3

Yield stress, maximum stress and strain-hardening exponents, n (Hollomon)

and n0 (Ludwik) obtained from the analyses of engineering and true stresse

true strain plots

Test condition Yield stress

(MPa)

Maximum

stress (MPa)

n n0

10�3 s�1, 1300 �C 862 883 0.12 0.59

10�4 s�1, 1100 �C 981 1177 0.62 0.98

10�4 s�1, 1200 �C 536 700 0.32 0.41

10�4 s�1, 1300 �C 489 542 0.27 0.41

10�4 s�1, 1350 �C 280 340 0.32 0.41

1470 R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

to have a slightly lower melting point, compared to that ofpure Mo. Hence, the self-diffusivity in a-Mo is expected tobe significantly large at 1200 �C (0.51Tm of Mo) or above, en-abling large-scale thermally activated plastic deformation ofthe bcc-structured a-Mo phase. Thermally activated deforma-tion is routinely observed in metals and alloys having bccstructure, characterized by relatively loose atomic packing,higher self-diffusivity and high Peierls barrier at lower temper-atures. The reduction in the yield strength of the MSB alloycould also be due to the softening of SiO2 particles, whichcomprise a significant volume fraction.

The true stressetrue strain data have been analyzed to cal-culate the strain-hardening exponent from equations proposedby Hollomon (Eq. (5)) [47] and Ludwik (Eq. (6)) [48]. TheHollomon equation states:

s¼ K3n ð5Þ

where K is a constant and n is the Hollomon strain-hardeningexponent. The Ludwik equation appears as:

s¼ s0þK03n0

p ð6Þ

where s0 is the flow stress corresponding to 0% plastic strain,3p is the true plastic strain, K0 is a constant, and n0 is the Lud-wik strain-hardening exponent. The choice of true plasticstrain excludes the elastic contribution of machine parts tothe true strain. The Eq. (6) can be re-written as:

logðs� s0Þ ¼ log K0 þ n0 log 3p ð6aÞ

The Eqs. (6) and (6a) lead to an expression for the rate ofstrain hardening appearing as:

ds=d3p ¼ n0�ðs� s0Þ=3p

�ð6bÞ

The values of n and n0 obtained by analyses of true stressetruestrain data are summarized in Table 3. The value of n or n0 at

0.00 0.01 0.02 0.03 0.04 0.050

4

8

12

16

20

24

Str

ain

Har

den

ing

Rat

e (G

Pa)

Plastic Strain

1100°C

1200°C

1300°C

1350°C

Fig. 12. Variation of the rates of strain hardening, ds/d3p with plastic strain, 3p

in the temperature range of 1100e1350 �C.

1100 �C is as high as 0.62 or 0.98, respectively, suggestive ofvery high rate of strain hardening. On the other hand, thestrain-hardening exponents corresponding to tests conductedat higher temperatures are lower, and do not change much be-tween 1200 �C and 1350 �C. The strain-hardening exponentsobtained using the Hollomon equation are more than that forpure Mo (n¼ 0.13) and most bcc metals (n between 0.1 and0.2), in spite of the fact that it is a-Mo, which shows plasticdeformation and strain hardening. The higher values ofstrain-hardening exponents of the MSB alloy are probablydue to the constrained plastic flow inside the a-Mo phase,which amounts to a multiaxial state of deformation.

Fig. 12 shows the variation of the rate of strain hardening,ds/d3p with increasing plastic strain, 3p for compression testscarried out between 1100 �C and 1350 �C with a strain rateof 10�4 s�1. As expected, the rate of strain hardening de-creases with increasing plastic strain at a fixed temperature,and with increasing temperature at a particular strain, as is ex-pected for thermally activated flow.

4. Conclusions

Alloys having the compositions of 76Moe14Sie10B(MSB), 77Moe12Sie8Be3Al (MSB3Al) or 73.4Moe11.2Sie8.1Be7.3Al (MSB7.3Al) were processed throughthe reaction hot pressing process and characterized. The fol-lowing conclusions may be drawn from the study:

(1) The phases, a-Mo, Mo3Si, and Mo5SiB2 have been ob-served in the microstructures of MSB, MSB3Al andMSB7.3Al alloys. In addition, SiO2 has been found dis-persed in the MSB alloy. The Al is present in the MSB3Aland MSB7.3Al alloys, either in solid solution inside thea-Mo and Mo3Si phases, or as a-Al2O3 particles formedby reaction of Al with SiO2.

(2) The room temperature flexural strength of about 400 MPahas been recorded for the MSB alloy, which is signifi-cantly higher than that of monolithic MoSi2. The contribu-tion to flexural strength is derived from the hard phasesin the microstructure, as well as the constrained plastic de-formation of a-Mo. The non-uniform distribution andreduced volume fraction of a-Mo in comparison to thebrittle phases or coarse boride particles or porosities actingas flaws result in much lower flexural strength of theMSB7.3Al alloy.

(3) The average values of room temperature fracture tough-ness of the MSB, MSB3Al and MSB7.3Al alloys deter-mined through testing of SENB specimens lie between 5and 7 MPaOm, though those found through the indentationmethods are slightly higher. Non-catastrophic and gracefulfailure has been observed in some of the samples, depend-ing on the position of the notch. Indentation cracking be-haviour shows a reduced rate of increase in crack lengthswith increasing load, suggesting R-curve type behaviour,particularly in case of the MSB3Al and MSB7.3Al alloys.Toughening mechanisms such as crack arrest or bridgingby the a-Mo particles, and increased tortuosity of the

1471R. Mitra et al. / Intermetallics 14 (2006) 1461e1471

crack paths by deflection and branching have beenobserved.

(4) The high compressive yield strengths and rate of strainhardening observed for the MSB alloy suggest excellentstrength retention at elevated temperatures. The yieldstress and maximum stress decreases with increasing tem-perature, but rises with increasing strain rate, suggestingoperation of thermally activated plastic flow. The strain-hardening exponents from the Hollomon and Ludwikequations are the highest at 1100 �C, and are less butwithout much variation between 1200 �C and 1350 �C.The rates of strain hardening decrease with increasingtemperature.

Acknowledgements

The authors are grateful to the technical staff in the Depart-ment of Metallurgical and Materials Engineering for miscella-neous assistance, Prof. P. K. Sinha, Dept. of MechanicalEngineering, IIT, Kharagpur and Dr. Y. R. Mahajan, AdvancedResearch Centre International, Hyderabad for extendingthe EDM facilities to prepare specimens. Thanks are alsodue to Mr. M. M. Srinivasa Rao, Technical Officer and Mr.V. V. Rama Rao, Scientist, Defence Metallurgical ResearchLaboratory (DMRL), Hyderabad for enthusiastic technical as-sistance. The authors also gratefully acknowledge the infra-structure support for part of this research obtained from Dr.J. Subrahmaniam, Group Head, Ceramics and CompositesGroup, DMRL and Dr. D. Banerjee, Chief Controller, Re-search and Development, Defence Research and Development,New Delhi.

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