ambient- to elevated-temperature fracture and fatigue ...ambient- to elevated-temperature fracture...

Ambient- to Elevated-Temperature Fracture and FatigueProperties of Mo-Si-B Alloys Role of Microstructure

JJ KRUZIC JH SCHNEIBEL and RO RITCHIE

Ambient- to elevated-temperature fracture and fatigue-crack growth results are presented for five Mo-Mo3Si-Mo5SiB2ndashcontaining -Mo matrix (17 to 49 vol pct) alloys which are compared to results forintermetallic-matrix alloys with similar compositions By increasing the -Mo volume fraction ductilityor microstructural coarseness or by using a continuous -Mo matrix it was found that improved frac-ture and fatigue properties are achieved by promoting the active toughening mechanisms specificallycrack trapping and crack bridging by the -Mo phase Crack-initiation fracture toughness valuesincreased from 5 to 12 MPa with increasing -Mo content from 17 to 49 vol pct and fracturetoughness values rose with crack extension ranging from 85 to 21 MPa at ambient temperaturesFatigue thresholds benefited similarly from more -Mo phase and the fracture and fatigue resistancewas improved for all alloys tested at 1300 degC the latter effects being attributed to improved ductilityof the -Mo phase at elevated temperatures

1m1m

I INTRODUCTION

A major limitation in the technological advancement ofaerospace engines and power generation is the lack of higher-temperature structural materials to replace conventionalnickel-base superalloys Indeed single-crystal nickel-basesuperalloys have essentially reached their temperature limitand are unsuitable for structural use above 1100 degC[1]

Unfortunately using higher melting-point (2000 degC) mate-rials presents several obstacles as adequate fracture fatigueoxidation and creep resistance must be maintained Forexample materials based on refractory metals such as molyb-denum typically suffer from poor oxidation and creep resis-tance While the formation of molybdenum silicides andborosilicides significantly improves these properties[234]

such intermetallic compounds are invariably brittle andprovide little resistance to fracture

To meet the need of developing useful structural molyb-denum-based alloys two distinct multiphase Mo-Si-B alloysystems have been proposed based on the phases (1) Mo3SiMo5SiB2 (T2) and Mo5Si3 (T1) by Akinc and co-work-ers[2ndash5] and (2) -Mo Mo3Si and Mo5SiB2 (T2) by Berczik[67]

With regard to fracture resistance the latter alloys hold morepromise since they contain the relatively ductile -Mo phasewith Si and B in solid solution in addition to the brittleintermetallic phases Mo3Si and Mo5SiB2

The microstructural arrangement of these phases is criticalin determining the properties Indeed while alloys containingmolybdenum particles surrounded by the hard but brittleintermetallic phases display definitive improvements intoughness compared to monolithic silicides[89] larger

improvements in fracture and fatigue resistance have beenachieved using a continuous -Mo matrix[10] Although these-Mo matrix alloys resemble nickel-base () superalloysin their similarly high volume fraction of intermetallic phasean additional difficulty for Mo-Si-B materials is that unlikethe Ni3Al phase Mo3Si and Mo5SiB2 exhibit essentiallyno ductility at ambient temperatures These factors suggestthat making the most effective use of the relatively ductilemolybdenum phase is important for achieving high fractureand fatigue resistance However since the development ofserviceable Mo-Si-B alloys will depend on compromisesbetween several material properties[11] there is a need tounderstand more specifically how different microstructuralparameters affect the fracture and fatigue behavior Accord-ingly it is the objective of this work to present results onthe fracture toughness and fatigue-crack propagation prop-erties at ambient to elevated (1300 degC) temperatures for aseries of five Mo-Si-B alloys with continuous -Mo matri-ces such results are compared to those for intermetallic-matrix alloys with similar compositions[8911] in an effort tounderstand the specific role of microstructure in influenc-ing the salient micromechanisms responsible for their frac-ture and fatigue behavior

II MATERIALS AND PROCEDURES

To generate microstructures with a continuous -Mophase powders were ground from arc-cast ingots with thecomposition Mo-20Si-10B at pct (Mo3Si-Mo5SiB2) andvacuum annealed to remove silicon from the surface and toleave an -Mo coating on each particle[12] The surface-modified powders were then hot-isostatically pressed inevacuated Nb cans for 4 hours at 1600 degC at 200 MPa pres-sure Five alloys were produced differing in volume frac-tion of the -Mo matrix and coarseness of the intermetallicparticles Polished metallurgical sections were used to char-acterize the microstructures and to determine the -Mo con-tent of each alloy These alloys contained Mo3Si (cubic A15structure) and Mo5SiB2 (tetragonal D81 structure) inter-metallic phases within a continuous -Mo matrix two alloys

METALLURGICAL AND MATERIALS TRANSACTIONS A US GOVERNMENT WORK VOLUME 36A SEPTEMBER 2005mdash2393NOT PROTECTED BY US COPYRIGHT

JJ KRUZIC Assistant Professor is with the Department of MechanicalEngineering Oregon State University Corvallis OR 97331 Contact e-mailjamiekruzicoregonstateedu JH SCHNEIBEL Senior Research StaffMember is with Oak Ridge National Laboratory Metals and CeramicsDivision Oak Ridge TN 37831 RO RITCHIE Professor is with theMaterials Sciences Division Lawrence Berkeley National Laboratoryand the Department of Materials Science and Engineering University ofCalifornia Berkeley CA 94720

Manuscript submitted January 3 2005

2394mdashVOLUME 36A SEPTEMBER 2005 METALLURGICAL AND MATERIALS TRANSACTIONS A

designated as F34 and M34 had 34 vol pct -Mo with initial

The alloy designations constitute a letter followed by a number Theletter refers to the relative coarseness of the microstructure (F fine M medium and C coarse) the number refers to the volume percent of the-Mo phase Fine medium and coarse correspond to 45 45 to 90 mand 90 to 180 m starting powders respectively

powder sizes of 45 and 45 to 90 m respectively whilethree coarse microstructured alloys had 17 46 and 49 volpct -Mo (designated C17 C46 and C49 respectively) andan initial powder size of 90 to 180 m (Figure 1) Notethat alloys F34 M34 and C49 correspond to the alloys des-ignated ldquofinerdquo ldquomediumrdquo and ldquocoarserdquo in Reference 10where initial fracture and fatigue data were presented Inaddition to the equilibrium Mo3Si and Mo5SiB2 phases smallamounts of the nonequilibrium Mo2B phase were detectedby X-ray diffraction

The chemical compositions were different from the ini-tial starting powders due to the vacuum annealing process

Final alloy compositions were determined by a combinationof inductively coupled-plasma spectroscopy and combustion-infrared absorption the latter technique was used to deter-mine the oxygen carbon and sulfur content

Resistance-curve (R-curve) fracture-toughness experimentswere performed on disk-shaped compact-tension DC(T)specimens (14-mm width 3-mm thickness) in general accor-dance with ASTM Standard E561[13] Specimens were pre-cracked by fatigue loading (fatigue conditions describedsubsequently) half-chevron notched specimens at room tem-perature until precracks had propagated beyond the half-chevron notch region (which was 300 to 600 m in length)Samples were then loaded monotonically in displacementcontrol at 1 mmin until the onset of crack extension At25 degC periodic unloads (10 to 20 pct of peak load) wereperformed to measure the unloading back-face strain com-pliance which was used to determine the crack length[14]

Elevated temperature tests at 1300 degC were conducted in agettered argon environment using (direct-current) electrical

Fig 1mdashOptical micrographs of the (a) F34 (b) M34 (c) C17 (d ) C46 and (e) C49 alloys showing continuous -Mo matrices with Mo3Si and T2 (etchedwith Murakamirsquos reagent)

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 36A SEPTEMBER 2005mdash2395

potential-drop techniques to monitor crack length Such ultra-high temperature testing is complex to perform full detailsof our procedures are described in Reference 15 At both25 and 1300 degC fracture resistance was assessed as a func-tion of crack length based on the loads associated with eachcrack length determined by the compliance or electric poten-tial methods In specific instances testing was paused toobserve crack profiles such that interactions of the crackwith microstructural features could be documented usingoptical microscopy and scanning electron microscopy (SEM)The R-curve tests were continued until catastrophic failureof the specimens occurred or more than at least 3 mm ofcrack extension was achieved at which point the test wasconcluded For cases where tests were interrupted or cata-strophic failure did not occur crack lengths determined bycompliance or electrical potential methods were verifiedusing optical measurements Since both compliance andpotential-drop techniques can be prone to errors due to crackbridging between the crack faces any discrepancies in cracklength were accounted for by assuming that they accumu-lated linearly with crack extension

Fatigue-crack growth testing (25 Hz sinusoidal waveform)was performed at 25 and 1300 degC in identical environments toR-curve testing in general accordance with ASTM standardE647[16] using computer-controlled servohydraulic testingmachines at a load ratio R (ratio of minimum to maximumloads) of 01 Crack-growth rates dadN were determined asa function of the stress-intensity range K using continuousload shedding to maintain a K gradient ([1K]dKda) of008 mm1 corresponding to increasing or decreasing Kconditions respectively The KTH fatigue thresholds opera-tionally defined at a minimum growth rate of 1010 to 1011

mcycle were approached under decreasing K conditionsTesting was again periodically interrupted to conduct opticalmicroscopy and SEM such that any errors in crack-length mea-surement could be accounted for as described previouslyCharacterization of the fatigue and fracture surfaces was per-formed using SEM after specimen failure

III RESULTS

A Alloy Compositions

Final compositions of the alloys are shown in Table I Asexpected the alloys with the highest -Mo content had thelowest Si concentration as a result of the vacuum anneal-ing process Unexpectedly the relatively low B concentra-tions for alloys C46 and C49 suggest that not only Si butalso B was partially removed during the vacuum annealing

step Nearly identical concentrations of Mo Si and B werefound for alloys F34 and M34 reflective of their nearly iden-tical -Mo volume fractions Alloys C46 and C49 howeverhad quite dissimilar measured compositions despite similar-Mo volume fractions the reasons for this are not clearUsing the chemical composition data from Table I and theequilibrium 1600 degC Mo-Si-B ternary phase diagram[17] esti-mates of the phase volume fractions were made based onthe lattice parameters for pure Mo and stoichiometric Mo3Siand Mo5SiB2 (Table II) Despite the simplifying assump-tions and the fact that nonequilibrium MoB2 was ignoredthere was relatively good agreement between the estimatesand the -Mo volume fractions measured from micrographsfor all alloys except C49 suggesting that the chemical analy-sis for alloy C49 may be in error

B Fracture Toughness

Room-temperature R-curves for the five Mo-Mo3Si-T2 alloysplotted in terms of the stress intensity K clearly indicate ris-ing fracture toughness with crack extension (Figure 2(a)) Themeasured R-curves are mostly linear although the catastrophicfailure of some samples suggests a sharp flattening out of theR-curves which was not measured The R-curve tougheningeffects were most significant for the coarse alloys C46 andC49 which had peak room-temperature toughnesses in excessof 20 MPa ie up to 7 times higher than that of monolithicmolybdenum silicides (MoSi2 Mo3Si) which have toughnessesof 3 to 4 MPa [1819] The current alloys are also tougherthan Mo-Si-B alloys with discontinuous -Mo phase (22 to38 vol pct) which displayed modest R-curve behavior withmuch lower peak toughness (40 to 75 MPa )[89]

Additionally experiments at 1300 degC on alloys M34 C17and C46 indicated that the fracture toughness improves athigher temperatures The R-curve data at 1300 degC are shownin Figure 2(b) along with room-temperature data for com-parison Increasing the temperature had a particularly signifi-cant effect on the crack-initiation toughness Ki which definesthe start of the R-curve specifically Ki values for alloys M34and C17 were respectively 63 and 65 pct higher at 1300 degCthan at 25 degC Again the toughness properties were mostimpressive for the coarse high -Mo alloy C46 which expe-rienced toughening at 1300 degC to such a degree that large-scalecrack blunting and deformation occurred and linear-elasticfracture mechanics (LEFM) was no longer a valid method forassessing the toughness using the present specimen sizeFigure 3(a) illustrates this behavior indeed a deformation zonegreater than 15 mm in size may be seen which is on the orderof sample dimensions and violates the small-scale yieldingrequirements for LEFM Crack blunting was also observed toa much smaller degree in the other alloys tested at 1300 degCas seen in Figure 3(b) for alloy M34

1m

1m

1m

Table I Compositions of Mo-Si-B Alloys (Atomic Percent)

Element F34 M34 C17 C46 C49

Mo 7668 7728 7362 8047 8514Si 1132 1132 1502 946 668B 1142 1092 1099 989 805Al 032 026 017 012 009Fe 011 001 008 004 003Ni mdash 013 mdash mdash mdashO 008 0005 005 001 001C 007 007 006 007 007S 0005 0007 0007 0005 0005

Table II Estimated Volume Percent of Each Phase Basedon Compositions in Table I

Phase F34 M34 C17 C46 C49

-Mo 375 375 183 491 665Mo3Si 219 219 403 146 47Mo5SiB2 406 406 414 363 288

Inconsistent with measured -Mo phase volume fraction


Fig 2mdashR-curves showing the fracture resistance of the continuous -Mo matrix Mo-Si-B alloys at (a) room temperature (open symbols) and (b) 1300 degC(closed symbols) Also shown are previously reported room-temperature values for unreinforced molybdenum silicides[1819]

Fig 3mdashOptical micrographs illustrating the crack blunting and damagezone ahead of the crack tips after R-curve testing at 1300 degC for alloys(a) C46 and (b) M34 Notice the order of magnitude difference in scalebetween the two photos Furthermore crack bridging is seen for alloy M34in (b) Nomarski differential image contrast was used for both imagesand the nominal crack growth direction was left to right

For alloy C46 the J-integral may be estimated based on themeasured crack-tip opening displacement by the relation[20]

[1]J dn s0 d

where 130 is the yield strength and dn is a dimensionlessfactor that varies between 03 and 1 depending upon the yieldstrain strain-hardening coefficient and stress state Limitedtensile tests at 1200 degC on alloy C49 which is similar inmicrostructure and composition to alloy C46 gave yield andultimate tensile strengths of 336 and 354 MPa respectively[10]

Thus using dn 1 which reflects a plane-stress conditionwith minimal strain hardening[20] a yield strength of 130 336 MPa and the measured of 10 m gives an estimateof the crack-initiation toughness of Ji 3360 Jm2 Whenlinear-elastic conditions prevail such a J value can be relatedto the stress intensity K under mode I conditions by

[2]

where E is the elastic modulus This allows an estimate ofthe plane-stress initiation toughness in terms of K whichshould reflect the value that would have been measured ifit were possible to use a larger specimen size to maintainsmall-scale yielding conditions From Eq[2] a plane-stressinitiation toughness of Ki 35 MPa was deduced usinga modulus E 360 GPa which is based on a simple rule ofmixtures using the volume fraction of the -Mo phase (E 325 GPa[21]) and the intermetallic phases (E 390 GPa[2223])The measured or deduced initiation toughness values foreach alloy are tabulated in Table III

C Fatigue-Crack Growth

Fatigue-crack growth results for the five alloys are shownin Figure 4 and may be characterized in terms of a classi-cal power-law relationship

[3]

It is apparent that at 25 degC the power-law exponents m areextremely high exceeding 78 for all alloys (Table III) whichis characteristic of brittle materials such as ceramics Fatigue-crack growth data are also shown for alloys M34 and C49at 1300 degC Although alloy M34 had a similarly high Kdependence the coarse high -Mo alloy C49 displayed atransition to more ductile fatigue behavior at 1300 degC withmore than an order of magnitude decrease in the power-lawexponent from 78 to 4 An exponent of 4 is similar to what

dagtdN CK m

1m

K 1JE


Table III Fracture and Fatigue Properties of Mo-Si-BAlloys with -Mo Matrix

Initiation Initiation Fatigue Toughness Toughness Threshold Power-LawKi at 25 degC Ki at 1300 degC KmaxTH at Exponent

Alloy (MPa ) (MPa ) 25 degC (MPa ) m at 25 degC

F34 80 NA 68 125M34 75 126 72 87C17 50 83 56 152C46 98 35 99 88C49 120 NA 106 78

Plane stress value estimated from measured crack tip opening displacement(CTOD)

NAmdashnot applicable

1m1m1m

Fig 4mdashFatigue-crack growth rate dadN data plotted as a function ofapplied stress intensity range K for continuous -Mo matrix Mo-Si-Balloys at room (25 degC) and elevated (1300 degC) temperatures The fatiguethreshold clearly increases with increasing -Mo vol pct Also note thedramatic decrease in power-law exponent (78 to 4) for alloy C49 whenthe temperature is raised from 25 degC to 1300 degC

is expected for ductile metals which typically have m val-ues in the range 2 to 4 for the midrange of growth rates[24]

The KTH thresholds ranged between 5 and 95 MPafor the five alloys at 25 degC (Figure 4) Specific values foreach alloy are given in terms of KmaxTH the maximum stressintensity at the fatigue threshold in Table III Here KmaxTH

was chosen as the most appropriate parameter since the fatigueof brittle materials is generally most influenced by Kmax witha lesser dependence on K[24] Note both expressions ofthe fatigue threshold are related however via the load ratio

[4]

At 1300 degC due to experimental difficulties and limited num-bers of samples data were not collected near the operationallydefined fatigue threshold however based on extrapolationof the data in Figure 4 the threshold for M34 is expectedto be similar to that at 25 degC whereas data for alloy C49suggest a decrease in the fatigue threshold at 1300 degC

KTH Kmax TH (1R)

1m

D Toughening Mechanisms

Observations of the crack profiles revealed crack trappingand crack bridging as the primary toughening mechanisms(Figure 5) Crack trapping refers to the local arrest of crackgrowth at specific microstructural features in this case the-Mo phase (Figure 5(a)) For the ldquotrappedrdquo crack to over-come this barrier higher applied driving forces are neededresulting in higher overall toughness for the material Cracktrapping is related to the inherent local resistance to crackingof the material and can be classified as an intrinsic tougheningmechanism that acts to raise the initiation toughness Ki

[24]

Conversely rising R-curve behavior or crack-growth tough-ness is indicative of extrinsic toughening which acts awayfrom the crack tip and develops with crack extension leadingto rising fracture resistance[24] Crack bridging which is observedin the present alloys (Figures 3(b) and 5) is one such mecha-nism whereby bridges of material span the wake of the crackresisting crack opening and sustaining some of the applied loadthat would otherwise contribute to crack advance Bridges form

Fig 5mdash(a) Crack trapping and bridging at the -Mo phase in alloy M34The crack locally arrests at the -Mo phase leaving -Mo bridges in thecrack wake (b) Crack bridging 33 mm behind the crack tip in alloy C46illustrating how this mechanism can account for the rising R-curve behaviorover several millimetres as seen in Fig 2 Both specimens were tested atroom temperature and etched with Murakamirsquos reagent


during crack extension leading to the rising toughness withcrack growth (ie rising R-curve behavior) Although inter-metallic bridges were occasionally observed (eg Figure 6(a))most bridges were composed of the -Mo phase and appearedto result from cracks renucleating in the Mo3Si and T2 phasesahead of the crack trapping points ie as the crack extendedcrack trapping at the -Mo acted to promote the formation ofcrack bridges (Figure 5(a)) The existence of bridges persistedover many millimeters behind the crack tip in alloys C17 C46and C49 (Figure 5(b)) which accounts for the rising toughnessseen in these alloys over similar amounts of crack extension(Figure 2) Bridges were observed to degrade during fatigue-crack growth however as shown in Figure 6

The increased fracture resistance at 1300 degC was associatedwith enhanced crack blunting (Figure 3) while extensive defor-mation of bridges in the crack wake was also seen (Figure 7)This observed behavior is indicative of the improved -Moductility at elevated temperatures Conversely the room-tem-perature fast fracture surfaces revealed some degree of inter-granular failure in the -Mo phase (Figure 8) which can beattributed to limited ductility of the -Mo at lower temperatures

Fig 6mdash(a) A bridge in alloy M34 located 400 m behind the tip of afatigue crack at room temperature (b) After 800 m of additional crackadvance the bridge has been destroyed The specimen was etched withMurakamirsquos reagent

Fig 7mdash-Mo bridges behind the crack tip for alloy M34 after R-curvetesting at 1300 degC Note the large amounts of ductility for the bridges at1300 degC relative to those tested at room temperature (Fig 5)

Fig 8mdashRoom-temperature fast fracture surfaces for alloys (a) M34 and(b) C49 Flat regions correspond to the cleavage fracture of the intermetallicphases while arrows point out regions where intergranular failure is appar-ent in the -Mo phase


IV DISCUSSION

A Role of -Mo Volume Fraction

It is apparent that the fracture toughness and fatigue-crackgrowth properties of the -Mo matrix alloys are progres-sively improved with increasing -Mo volume fraction asshown in Figure 9 where the Ki initiation toughness andKmaxTH fatigue threshold values from Table III are includedAlso shown in Figure 9 are results on intermetallic matrixMo-Si-B alloys with similar compositions (22 to 38 vol pct-Mo)[89] which exhibit improved properties with larger-Mo volume fractions as well The linear relations on thesemilog plot in Figure 9 indicate that Ki and KmaxTH increaseexponentially (fitted lines) with the -Mo volume fractionat room temperature for both of these classes of Mo-Si-Bmaterials however this trend clearly cannot continue to100 pct -Mo as all alloys would have the same toughnessat this limit

Steady-state toughness values ie the peak values fromthe plateau region of the R-curve were not measured dueto inadvertent specimen failure or limited specimen size (ielarge-scale bridging conditions) and accordingly are notcompared in Figure 9 Based on Figure 2 however it appearsthat the crack-growth toughness or rising part of the R-curvedoes benefit from higher levels of -Mo This may be con-cluded from Figure 2(a) when one considers the three alloysC17 C46 and C49 which all were processed from the samecoarse starting powder yet had different final -Mo contentsWhen comparing their R-curve behavior it is seen that thealloys with more -Mo phase have R-curves that rise moresteeply This would imply that peak toughnesses increasewith increasing -Mo by an amount larger than the differ-ences in Ki

These trends may be understood by considering the salienttoughening mechanisms in these alloys which have been

Fig 9mdashComparison of the crack-initiation toughness and fatigue thresh-old values for the present alloys (Table III) with those for intermetallicmatrix alloys that had similar compositions from Refs 8 and 9 all valuesare plotted as a function of -Mo volume percent

identified as crack trapping and crack bridging The rela-tively ductile -Mo serves as the trapping phase in boththe -Mo and intermetallic matrix alloys which is fully con-sistent with the observed increase in crack-initiation tough-ness with increasing -Mo volume fraction (Figure 9)Additionally the potency of crack bridging in the presentalloys is aided by increasing the amount of -Mo Sincemost bridging is by the -Mo phase as the -Mo contentincreases it is reasoned that more or larger bridges developas the crack advances leading to the observed gains in extrin-sic toughness

For fatigue-crack growth behavior room-temperatureKmaxTH thresholds (R 01) also rose with increasing -Mocontent and were found to be essentially equal to Ki (exceptat 49 pct -Mo where Ki slightly exceeded KmaxTH) (Fig-ure 9) This suggests that the intrinsic mechanisms for crackadvance at ambient temperature are identical under cyclicand monotonic loading for the alloys with -Mo content46 pct This behavior is common in brittle materials wherehigh power-law exponents are observed and fatigue-crackgrowth is typically a result of a cyclic-loading induced degra-dation in the extrinsic toughening mechanisms[24] In thiscase the degradation occurs by the destruction of bridgesdue to cyclic loading as seen in Figure 6 for alloy M34Even for the 49 pct -Mo alloy the power-law exponentm remains extremely high 78 implying that fatigue-crackadvance is still dominated by brittle mechanisms Indeedductile fatigue mechanisms are associated with m 2 to 4while Ki is often an order of magnitude or more larger thanKmaxTH for ductile metals[24]

One advantage of the brittle room-temperature fatiguebehavior of the present alloys is the high fatigue thresholdsrelative to traditional (ductile) metallic alloys where thresh-old values are typically below 5 MPa For high-frequencyfatigue applications where the fatigue-crack growth portionof the lifetime is insignificantly small high fatigue thresh-olds are advantageous since higher stresses can be achievedwithout causing crack advance However the accompany-ing large power-law exponents are not desirable for fatigueapplications where a degree of damage tolerance is requiredsince once a crack starts growing failure will quickly fol-low unless the stresses are decaying fast enough to causecrack arrest

B Role of -Mo Ductility

Results at 1300 degC reveal the effect of -Mo ductility onthe fracture and fatigue properties of these alloys Indeedthe main effect of elevated temperature is thought to be theldquoductilizationrdquo of the -Mo phase while the Mo3Si and T2phases remain brittle The initiation toughness valuesincreased by 65 pct for the C17 and M34 alloys while esti-mated gains for alloy C46 were on the order of 257 pctthough this latter value is expected to be somewhat elevateddue to the plane-stress conditions at 1300 degC (Figure 9) Itis interesting to note that at 1300 degC the initiation tough-ness of alloys C17 and M34 exceeded the room-tempera-ture Ki values for alloys with higher -Mo volume fractioneg 34 and 49 pct respectively This implies that less -Mowould be needed to achieve a given room-temperature tough-ness provided the ductility is improved This is importantbecause low -Mo volume fractions are desirable from the

1m


standpoint of improved oxidation and creep resistance inthese alloys[1125ndash28]

Figure 8 shows SEM micrographs of room-temperaturefast fracture surfaces for alloys M34 and C49 In these andindeed the other alloys there is clear evidence of inter-granular fracture of the -Mo phase indicative of the lowroom-temperature ductility of the -Mo Intergranular fail-ure in molybdenum is typically associated with impuritiessuch as oxygen which is known to segregate to and embrit-tle the grain boundaries[29] Carbon is thought to offset thiseffect to some degree with considerable ductility achievedin unalloyed molybdenum when the CO ratio is greater thantwo[30] although it is not known if similar effects may beexpected in the present -Mo solid solution Alloys F34 andC17 both had CO ratios close to unity while alloy M34had a high CO ratio (14) the C and O distribution amongthe phases however is unknown (Table I) For alloys C49and C46 the C concentration was below the detectabilitylimit making an accurate assessment of the CO ratio impos-sible in these alloys Observations of intergranular failurefor alloy M34 (Figure 8(a)) which had a CO ratio of 14suggest that other impurities may play an important roleeither by direct segregation to the grain boundaries or byaffecting the interaction of carbon and oxygen in the -Mosolid solution making the critical CO ratio to achieve ade-quate ductility different than expected for unalloyed molyb-denum Additionally Si present in the -Mo solid solutionis known to increase the hardness significantly[31] whichlikely reduces the ductility this suggests that minimizingthe Si in the -Mo phase may be another way to improvethe -Mo ductility Thus careful compositional control ofthe -Mo phase is expected to be an important factor inimproving the fracture resistance of Mo-Si-B alloys

Although the enhanced -Mo ductility at higher temper-atures has a marked influence on the ductility of the bridges(Figure 7) the crack-growth toughness was relatively unaf-fected This may be understood by considering that theincrease in ductility of Mo with increasing temperature isaccompanied by a decrease in the strength The enhancedtoughness due to ductile-phase bridging may be expressedin terms of the strain-energy release rate G by[323334]

The strain-energy release rate is an alternative measure of toughness thatmay be related to the stress intensity for a linear elastic material by

where E is the appropriate elastic mod-ulus (E in plane stress E(1 2) in plane strain where is Poissonrsquosratio) and is the shear modulus

[5]

where t is the size of the bridges 130 is the yield strengthVf is the volume fraction of bridges and is a work ofrupture parameter that depends on the strain hardening duc-tility and plastic constraint of the ductile phase Althoughrising ductility is beneficial to the bridging contributionthrough the parameter this will be offset by lower yieldstrengths at 1300 degC However in general these factors willnot completely offset and there is potential for improvingthe room-temperature bridging contribution by increasingthe ductility provided the yield strength remains high enoughto realize this benefit

With respect to fatigue the behavior of alloy C49 changeddramatically at 1300 degC with the power-law exponent fallingto 4 characteristic of ductile metals The lower exponent

G Vf s0 tx

G K 2I gtEiquest K 2

II gtEiquest K 2III gt2m

allows for more damage tolerance in the material as a con-siderable amount of the fatigue life will be spent growingan introduced flaw large enough to cause catastrophic failure

C Role of -Mo Matrix

The present alloys all had continuous -Mo matrices andtheir properties are compared in Figure 9 to data from pre-vious studies which looked at intermetallic matrix Mo-Si-B alloys of similar composition[89] At a given -Movolume fraction the fracture and fatigue properties wereimproved for the -Mo matrix materials particularly at1300 degC This behavior is attributed to higher effectivenessof the crack trapping and bridging mechanisms when thereis a continuous -Mo matrix since the crack cannot avoidthe relatively ductile phase Gains at room temperaturewere not spectacular however these improve with increas-ing -Mo volume fraction Indeed at 22 pct -Mo thebenefit to Ki is roughly 07 MPa or 15 pct while at 38pct -Mo the gain is 17 MPa or 26 pct This effectis enhanced when the -Mo ductility is higher achievedhere by testing at 1300 degC For 22 pct -Mo the -Momatrix alloys have higher Ki values by 08 MPa or10 pct while at 38 pct -Mo the increase is 61 MPa or 62 pct Thus using an -Mo matrix material is morebeneficial at higher -Mo fractions and ductility Since the-Mo matrix is thought to enhance toughness by forcingthe crack to interact with the -Mo phase it is not sur-prising that larger amounts and more ductile -Mo elevatesthis effect

D Role of Microstructural Scale

Comparing alloys M34 and F34 which both had 34 pct-Mo phase with different starting powder sizes the micro-structural size scale appeared to have no effect on the intrinsictoughness or fatigue threshold at room temperature (Table III)There was an effect on the extrinsic toughening (bridging)however as the R-curve rose higher for alloy M34 (Figure 2)Additionally crack stability was observed to be improvedin alloy M34 with far fewer samples breaking catastrophi-cally during precracking and testing reflective of easierbridge formation stabilizing crack propagation This pointis illustrated further by the behavior of alloy C17 for whichstable rising R-curve behavior was readily achieved overseveral millimeters despite the low amount (17 pct) of -Mophase Alloy C17 had the largest starting powder size andthus had a coarser microstructural scale compared to alloysF34 and M34 (Figure 1) Toughening contributions fromcrack bridging are typically found to be enhanced withcoarser microstructures as seen in a variety of intermetallicscomposites and ceramics[353637]

E Comparing to the Nb-Si System

The Nb-Si system is another class of refractory metal-silicidendashbased materials that has received considerableattention for potential high-temperature gas-turbine appli-cations[38ndash42] This system presents similar challenges toMo-Si-B with the Nb solid solution phase (Nbss) beingductile but easily oxidized and the niobium silicides beingvery brittle with relatively good oxidation resistance Fracture

1m1m

1m1m


and fatigue properties have been reported for alloys with 20to 95 pct metallic Nbss these properties are compared toMo-Si-B alloys from the present study and References 8 and9 in Figure 10 When the crack-initiation toughness Kiis compared at 50 pct metallic phase alloy C49 (49 pct-Mo Ki 12 MPa ) is similar to the Nb-Ti-Hf-Cr-Al-Simetal and silicide composite (MASC) with 54 pct metallicphase where Ki 7 to 13 MPa [39] Furthermore extrap-olation of the initiation toughness data for Mo-Si-B alloys(Figure 10(a)) to higher -Mo content implies that Ki val-ues should be comparable to the best Nb-Sindashbased alloysalthough such high metallic volume fractions should seriouslycompromise the oxidation resistance of both systems Alsoshown in Figure 10(a) are Kc fracture toughness values (iecomputed from the peak load at failure) from a study onNb-Sindashbased multiphase alloys[42] The R-curves were notmeasured in that study since little R-curve behavior wasobserved accordingly these Kc values should approach theinitiation toughness Ki and be useful for comparison Thesealloys which often contained Laves phases as well as sili-cides exhibit slightly superior toughness at low metallic vol-ume percents but in general appear to fall in line with thepresent Mo-Si-B alloys as the metallic phase volume percentincreases

One shortcoming of the Mo-Si-B alloys investigated sofar is that the R-curves rise much more slowly than some ofthe Nb-Si systems investigated[3940] The latter can achievetoughnesses over 20 MPa after only a few hundreds ofmicrometers of growth[3940] instead of several millimeters forthe present alloys Enhanced extrinsic toughening is oftenachieved through microstructural modification (eg Refer-ences 36 and 43) and considering some Nb-Sindashbased alloysexhibit essentially no R-curve behavior[42] there is most likelyroom for improvement in the Mo-Si-B alloys For examplethe Nb-Sindashbased alloys described in References 39 and 40had highly anisotropic microstructures achieved by either

1m

1m

1m

directional solidification or extrusion processing with thefracture properties measured transverse to the extrusion orsolidification direction This most likely impacted the R-curvebehavior significantly producing similar microstructures forMo-Si-B alloys would appear to be a promising approach

Figure 10(b) compares the room-temperature fatiguethresholds of various Mo-Si-Bndash and Nb-Sindashbased alloys fewdifferences are apparent except for the fact that the Mo-Si-Balloys had lower metal phase content The reported power-law exponents however were as low as m 53[41] for someNb-Si alloys reflecting the higher metallic phase contentand possibly more ductility of the Nbss phase relative to the-Mo for the alloys examined Improving the -Mo ductil-ity by controlling its composition (eg Si content) and impu-rities (eg oxygen) should thus be a high priority for thedevelopment of Mo-Si-B alloys

V CONCLUSIONS

Based on an experimental investigation into the ambient-to high-temperature fracture toughness and fatigue-crack prop-agation behavior of five Mo-Si-B alloys containing Mo3Siand Mo5SiB2 intermetallic phases dispersed within a contin-uous -Mo matrix the following conclusions may be made

1 The -Mo matrix Mo-Si-B alloys exhibit fracture andfatigue resistance that is far superior to that of unrein-forced silicides indeed fracture toughnesses in excessof 20 MPa may be achieved for -Mo volume frac-tions 45 pct

2 Crack trapping and crack bridging by the -Mo phase werefound to be the primary toughening mechanisms in thesealloys with the latter leading to rising fracture resistancewith crack extension Furthermore both of thesemechanisms benefited from higher -Mo volume fractionsleading to improved fracture and fatigue properties

1m

Fig 10mdash(a) No Crack-initiation toughness and (b) fatigue threshold values for Mo-Si-B alloys (present results and References 8 and 9) compared to pub-lished results for NB-Sindashbased materials[38ndash42] When assessed as a function of metallic-phase volume percent the Mo-Si-B alloys appear in line with mostNb-Si alloys in terms of both Ki and KmaxTH Note the toughness values from Ref 42 were Kc values which may elevate them relative to Ki but little R-curve behavior was reported making these comparisons reasonable


3 Fracture resistance was improved at 1300 degC indicatingthe important role that -Mo ductility plays in determiningthe mechanical properties of these alloys Based on theseresults it appears that a given level of fracture resistancemay be achieved with lower -Mo volume fraction byimproving -Mo ductility a desirable feature since -Mocompromises the oxidation and creep resistance

4 Using a continuous -Mo matrix instead of an inter-metallic matrix is also beneficial for the fracture andfatigue properties with the -Mo matrix having a largereffect with increasing -Mo content and ductility

5 Coarser microstructural size scales also benefit the fractureand fatigue behavior specifically by aiding bridging andcrack stability

6 When compared as a function of metallic-phase volumepercent the current Mo-Si-B alloys appear to be similarto most Nb-Sindashbased alloys with respect to their crack-initiation toughness and fatigue threshold values How-ever their power-law slopes are much larger than thoseof niobium silicides which implies poor tolerance tofatigue crack growth

7 The success of Mo-Si-B materials for turbine-blade andother high-temperature applications will depend on acompromise between several mechanical properties andthe oxidation resistance The present work defines therole of various microstructural variables in determiningthe fracture and fatigue properties and is intended to pro-vide guidelines for the design of superior alloys

ACKNOWLEDGMENTS

The mechanical property measurements in this work weresupported by the Office of Fossil Energy Advanced ResearchMaterials Program WBS Element LBNL-2 and themicrostructural analysis was supported by the Office of Sci-ence Office of Basic Energy Sciences Division of MaterialsSciences and Engineering United States Department ofEnergy under Contract No DE-AC03-76SF0098 with theLawrence Berkeley National Laboratory for ROR and JJKJHS acknowledges funding of the processing part of thiswork by the Office of Fossil Energy Advanced ResearchMaterials (ARM) Program WBS Element ORNL-2(I)United States Department of Energy under Contract NoDE-AC05-00OR22725 with Oak Ridge National Laboratorymanaged by UTndashBattelle

REFERENCES1 MJ Donachie Superalloys A Technical Guide 2nd ed ASM Metals

Park OH 2002 p 4392 MK Meyer MJ Kramer and M Akinc Intermetallics 1996 vol 4

pp 273-813 M Akinc MK Meyer MJ Kramer AJ Thom JJ Heubsch and

B Cook Mater Sci Eng 1999 vol A261 pp 16-234 MK Meyer and M Akinc J Am Ceram Soc 1996 vol 79

pp 2763-665 MK Meyer AJ Thom and M Akinc Intermetallics 1999 vol 7

pp 153-626 DM Berczik United Technologies Corporation East Hartford CT

US Patent 5595616 19977 DM Berczik United Technologies Corporation East Hartford CT

US Patent 5693156 1997

8 H Choe D Chen JH Schneibel and RO Ritchie Intermetallics2001 vol 9 pp 319-29

9 H Choe JH Schneibel and RO Ritchie Metall Mater Trans A2003 vol 34A pp 25-239

10 JJ Kruzic JH Schneibel and RO Ritchie Scripta Mater 2004vol 50 pp 459-64

11 JH Schneibel RO Ritchie JJ Kruzic and PF Tortorelli MetallMater Trans A 2005 vol 36A pp 525-31

12 JH Schneibel MJ Kramer and DS Easton Scripta Mater 2002vol 46 pp 217-21

13 ASTM E561-98 Annual Book of ASTM Standards vol 0301 Met-als-Mechanical Testing Elevated and Low-Temperature Tests Met-allography ASTM West Conshohocken PA 2002 pp 534-46

14 CJ Gilbert JM McNaney RH Dauskardt and RO Ritchie J Test-ing Eval 1994 vol 22 pp 117-20

15 D Chen CJ Gilbert and RO Ritchie J Testing Eval 2000 vol 28pp 236-41

16 ASTM E647-00 Annual Book of ASTM Standards vol 0301 Metals-Mechanical Testing Elevated and Low-temperature Tests Metallog-raphy ASTM West Conshohocken PA 2002 pp 595-635

17 DM Dimiduk and JH Perepezko MRS Bull 2003 vol 28 pp 639-4518 KT Venkateswara Rao WO Soboyejo and RO Ritchie Metall

Trans A 1992 vol 23A pp 2249-5719 I Rosales and JH Schneibel Intermetallics 2000 vol 8 pp 885-8920 CF Shih J Mech Phys Solids 1981 vol 29 pp 305-2621 A Buch Pure Metals Properties a Scientific-Technical Handbook

ASM INTERNATIONAL Materials Park OH 1999 p 30622 K Ito K Ihara K Tanaka M Fujikura and M Yamaguchi Inter-

metallics 2001 vol 9 pp 591-60223 JG Swadener I Rosales and JH Schneibel in High-Temperature

Ordered Intermetallic Alloys IX JH Schneibel S Hanada KJHemker RD Noebe and G Sauthoff eds Materials Research Soci-ety Warrendale PA 2001 pp N421-N426

24 RO Ritchie Int J Fract 1999 vol 100 pp 55-8325 JH Schneibel CT Liu DS Easton and CA Carmichael Mater

Sci Eng 1999 vol A261 pp 78-8326 JH Schneibel PF Tortorelli MJ Kramer AJ Thom JJ Kruzic

and RO Ritchie in Defect Properties and Related Phenomena inIntermetallic Alloys EP George MJ Mills H Inui and G Eggelereds Materials Research Society Warrendale PA 2003 pp 53-58

27 V Supatarawanich DR Johnson and CT Liu Mater Sci Eng2003 vol A344 pp 328-39

28 JH Schneibel Intermetallics 2003 vol 11 pp 625-3229 A Kumar and BL Eyre Proc R Soc London A 1980 vol 370

pp 431-5830 J Wadsworth TG Nieh and JJ Stephens Scripta Mater 1986

vol 20 pp 637-4231 NE Promisel The Science and Technology of Tungsten Tantalum

Molybdenum Niobium and Their Alloys Pergamon Press The MacMillanCompany New York NY 1964 p 588

32 LS Sigl PA Mataga BJ Dalgleish RM McMeeking and AGEvans Acta Metall 1988 vol 36 pp 945-53

33 MF Ashby FJ Blunt and M Bannister Acta Metall 1989 vol 37pp 1847-57

34 BD Flinn M Ruumlhle and AG Evans Acta Metall 1989 vol 37pp 3001-06

35 DR Bloyer KT Venkateswara Rao and RO Ritchie Metall MaterTrans A 1998 vol 29A pp 2483-96

36 JP Campbell KT Venkateswara Rao and RO Ritchie MetallMater Trans A 1999 vol 30A pp 563-77

37 RW Steinbrech A Reichl and W Schaarwaumlchter J Am CeramSoc 1990 vol 73 pp 2009-15

38 BP Bewlay MR Jackson J-C Zhao PR Subramanian MGMendiratta and JJ Lewandowski MRS Bull 2003 vol 29 pp 646-53

39 BP Bewlay MR Jackson and HA Lipsitt Metall Mater TransA 1996 vol 27A pp 3801-08

40 JD Rigney and JJ Lewandowski Metall Mater Trans A 1996vol 27A pp 3292-306

41 WA Zinsser Jr and JJ Lewandowski Metall Mater Trans A 1998vol 29A pp 1749-57

42 KS Chan and DL Davidson Metall Mater Trans A 2003 vol 34App 1833-49

43 CJ Gilbert JJ Cao WJ Moberly Chan LC De Jonghe and RORitchie Acta Metall Mater 1996 vol 44 pp 3199-214


designated as F34 and M34 had 34 vol pct -Mo with initial

The alloy designations constitute a letter followed by a number Theletter refers to the relative coarseness of the microstructure (F fine M medium and C coarse) the number refers to the volume percent of the-Mo phase Fine medium and coarse correspond to 45 45 to 90 mand 90 to 180 m starting powders respectively

powder sizes of 45 and 45 to 90 m respectively whilethree coarse microstructured alloys had 17 46 and 49 volpct -Mo (designated C17 C46 and C49 respectively) andan initial powder size of 90 to 180 m (Figure 1) Notethat alloys F34 M34 and C49 correspond to the alloys des-ignated ldquofinerdquo ldquomediumrdquo and ldquocoarserdquo in Reference 10where initial fracture and fatigue data were presented Inaddition to the equilibrium Mo3Si and Mo5SiB2 phases smallamounts of the nonequilibrium Mo2B phase were detectedby X-ray diffraction

The chemical compositions were different from the ini-tial starting powders due to the vacuum annealing process

Final alloy compositions were determined by a combinationof inductively coupled-plasma spectroscopy and combustion-infrared absorption the latter technique was used to deter-mine the oxygen carbon and sulfur content

Resistance-curve (R-curve) fracture-toughness experimentswere performed on disk-shaped compact-tension DC(T)specimens (14-mm width 3-mm thickness) in general accor-dance with ASTM Standard E561[13] Specimens were pre-cracked by fatigue loading (fatigue conditions describedsubsequently) half-chevron notched specimens at room tem-perature until precracks had propagated beyond the half-chevron notch region (which was 300 to 600 m in length)Samples were then loaded monotonically in displacementcontrol at 1 mmin until the onset of crack extension At25 degC periodic unloads (10 to 20 pct of peak load) wereperformed to measure the unloading back-face strain com-pliance which was used to determine the crack length[14]

Elevated temperature tests at 1300 degC were conducted in agettered argon environment using (direct-current) electrical

Fig 1mdashOptical micrographs of the (a) F34 (b) M34 (c) C17 (d ) C46 and (e) C49 alloys showing continuous -Mo matrices with Mo3Si and T2 (etchedwith Murakamirsquos reagent)





III RESULTS







1m

1m

1m






-Mo 375 375 183 491 665Mo3Si 219 219 403 146 47Mo5SiB2 406 406 414 363 288






[1]J dn s0 d



[2]




[3]


dagtdN CK m

1m

K 1JE





F34 80 NA 68 125M34 75 126 72 87C17 50 83 56 152C46 98 35 99 88C49 120 NA 106 78



1m1m1m





[4]


KTH Kmax TH (1R)

1m



[24]










IV DISCUSSION











1m







[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997







































III RESULTS







1m

1m

1m






-Mo 375 375 183 491 665Mo3Si 219 219 403 146 47Mo5SiB2 406 406 414 363 288






[1]J dn s0 d



[2]




[3]


dagtdN CK m

1m

K 1JE





F34 80 NA 68 125M34 75 126 72 87C17 50 83 56 152C46 98 35 99 88C49 120 NA 106 78



1m1m1m





[4]


KTH Kmax TH (1R)

1m



[24]










IV DISCUSSION











1m







[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997







































[1]J dn s0 d



[2]




[3]


dagtdN CK m

1m

K 1JE





F34 80 NA 68 125M34 75 126 72 87C17 50 83 56 152C46 98 35 99 88C49 120 NA 106 78



1m1m1m





[4]


KTH Kmax TH (1R)

1m



[24]










IV DISCUSSION











1m







[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997







































F34 80 NA 68 125M34 75 126 72 87C17 50 83 56 152C46 98 35 99 88C49 120 NA 106 78



1m1m1m





[4]


KTH Kmax TH (1R)

1m



[24]










IV DISCUSSION











1m







[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997










































IV DISCUSSION











1m







[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997









































[5]



G Vf s0 tx










1m1m

1m1m




1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997






































1m

1m

1m



V CONCLUSIONS




1m








ACKNOWLEDGMENTS









US Patent 5693156 1997









































ACKNOWLEDGMENTS









US Patent 5693156 1997



































ambient- to elevated-temperature fracture and fatigue ...ambient- to elevated-temperature fracture...

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