Cleavage fracture in B2 aluminides

Acta metalL mater. Vol. 40, No. 10, pp. 2727-2737, 1992

0956-7151/92 $5.00 + 0.00 Copyright c~,~1992 Pergamon Press Ltd

Printed m Great Britain. All rights reserved

CLEAVAGE FRACTURE IN B2 ALUMINIDES K.-M. C H A N G t, R. D A R O L I A 2 and H. A. L I P S I T T 3 ~GE Corporate Research and Development, Schenectady, NY 12301, ~-GE Aircraft Engines, Evendale, OH 45215 and 3Wright State University, Dayton, OH 45435, U.S.A. (Received 13 December 1991)

Abstract--Cleavage fracture of B2 atuminide single crystals, including FeA1, NiAI and CoAl, has been investigated at temperatures below their ductile-brittle transition temperatures. Single-edge notched bend specimens oriented along specific crystalline directions were tested by 4-point bending. The fracture resistance was highly anisotropic because of the existence of a preferred cleavage plane in these B2 aluminide crystals, With a deep through notch NiAI and CoAl crystals that have high ordering energies generally cleave on {110} planes, while substoichiometric FeA1 having a low ordering energy shows {I00} cleavage as do most b.c.c, metals. In the case of NiA1, a transition fracture region, composed of fracture facets on {511 } transient planes, appears at the initial cracking stage, followed by final cleavage on {110}. Different stoichiometric effects on the fracture toughness of B2 aluminides are observed when the A1 concentration is reduced. A general discussion on different mechanistic models has been used to explain the preferred cleavage planes in B2 structures. The intrinsic fracture toughness of an aluminide crystal can be determined by an ideal fracture test in which the cleavage plane is parallel to the notch plane and is normal to the applied stress. Because of geometric constraints an increased fracture resistance is obtained when the natural cleavage plane is not parallel to the notch plane, and the anisotropy of fracture toughness can be explained by a simple approach of resolved normal stress intensity. R6sum~---On 6tudie la rupture par clivage de monocristaux d'aluminiures de structure B2, FeA1, NiA1 et CoAl, ~ des temp6ratures situ6es au dessous de la transition ductile-fragile. Des 6chantillons fi entaille simple, orient6s dans des directions cristallographiques sp6cifiques, sont test6s en flexion sur quatre points. On trouve que la r6sistance fi la rupture est fortement anisotrope par suite de l'existence d'un plan de clivage pref6rential dans ces cristaux d'aluminiure de type B2. Avec une entaille profonde, les cristaux de NiAI et de CoAl, qui ont des ~nergies de raise en ordre ~lev6es, se clivent g~n~ralement dans des plans {110} comme la plupart des m6taux c.c. Dans le cas du NiA1, une region de transition de rupture, compos6e de facettes de rupture sur des plans transitoires {511}, apparait au stade initial de la fissuration, suivie du clivage final dans {110}. On observe diff6rents effets de stoechiom6trie sur la t6nacit6 fi la rupture des aluminiures de type B2 lorsque la concentration en AI est r6duite. On utilise une discussion g6n6rale sur les diff6rents mod61es de m6canismes pour expliquer les plans de clivage pr6f6rentiels dans les structures B2. La t~nacit6 intrins~que fi la rupture d'un cristal d'aluminiure peut &re d&ermin6e par un essai de rupture ideal dans lequel le plan de clivage est parall~le au plan d'entaille et est normal fi la contrainte appliqu6e. Par suite des contraintes g6om~triques, la resistance fi la rupture augmente lorsque le plan naturel de clivage n'est pas parall61e au plan d'entaille, et l'anisotropie de la t6nacite ~ la rupture peut s'expliquer simplement d'apr6s l'intensit6 de la contrainte normale r6duite.

Zusammenfassung--Der Spaltbruch von Einkristallen der B2-Aluminiden FeA1, NiA1 und CoAl wird im Temperaturbereich unterhalb des duktil-spr6den l~berganges untersucht. Einseitig gekerbte Biegeproben, die entlang bestimmter Kristallrichtungen orientiert sind, werden im Vierpunktbiegeversuch verformt. Der Bruchwiderstand ist hoch anisotrop, well diese B2-Aluminide eine bevorzugte Spaltebene aufweisen. Bei einer tiefen Kerbe spalten NiA1 und CoAl, die hohe Ordnungsenergie aufweisen, entlang von {ll0}-Ebenen, wohingegen unterst6chiometrisches FeAI mit niedriger Ordnungsenergie entlang von {100} spaltet wie die meisten k.r.z. Metalle. Im Falle des NiA1 erscheint ein 13bergangsbereich, in dem Bruchfacetten auf {511}-(3bergangsebenen zu Beginn des Spaltens auftreten, gefolgt vom Spalten auf letztlich {110}. Wird die AI-Konzentration erniedrigt, dann treten verschiedene St6chiometrie-Effekte bei der Bruchz/ihigkeit auf. Mit einer allgemneinen Diskussion verschiedener mechanistischer Modelle wird die in den B2-Strukturen bevorzugte Spaltebene erkl/irt. Die intrinsische Bruchfestigkeit eines Aluminidkristalles kann mit einem idealen Bruchtest bestimmt werden, bei dem die Spaltebene parallel zu einer Kerbebene und normal zur/iul3eren Spannung ist. Wegen der geometrischen Einschr/inkungen ergibt sich ein erh6hter Bruchwiderstand, wenn die natfirliche Spaltebene nicht parallel zur Kerbebene liegt; die Anisotropie der Bruchz/ihigkeit kann danach einfach anhand der Normalspannungsintensit/it erkl/irt werden.

INTRODUCTION M o s t body-centered-cubic (b.c.c.) metals, such as Fe, Mo, C r a n d W, etc., u n d e r g o brittle or cleavage

fracture at low temperatures. The ductile-to-brittle transition in Fe crystals a n d alloy steels has been o f great interest in b o t h scientific research a n d engineering applications [1, 2]. So far n o clear e x p l a n a t i o n

2727

2728

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

is available for predicting their fracture plane; the b.c.c, transition metals cleave on the { 100} planes in spite of the fact that the {110} planes are more closely packed with atoms [3]. The presence of impurity atoms in interstitial solid solution tends to increase the ductile-to-brittle transition temperature (DBTT) of b.c.c, metals. B2 aluminides are the intermetallic compounds that have an ordered b.c.c, structure (CsC1 type). Figure 1 shows the transition metals that can form binary B2 aluminides with aluminum. The preferred bonding between A1 and transition metal atoms causes each type of atom to locate at specific sites in the b.c.c, crystal lattice. Although the ordering energy varies from one aluminide phase to the other, these ordered B2 intermetallics have a higher DBTT than b.c.c, metals. A brittle fracture occurs in most B2 aluminides at room temperature. Early investigations on the mechanical behavior of B2 aluminides were limited to studies of the deformation character. Usually, compression tests were performed at temperatures of interest to avoid the premature failure associated wtih material brittleness. The active slip systems may be identified by employing slip trace analysis or by using transmission electron microscopy (TEM). Some summary reviews have recently appeared in the literature [4, 5]. However, very few studies on the fracture behavior of B2 aluminides have been reported. Because intergranular failure often occurs in polycrystalline materials, the identification of the ductile-to-brittle transition in B2 aluminides should be performed on single crystals. Pascoe and Newey have reported briefly that failure in stoichiometric NiA1 single crystals was generally by cleavage on {110} planes [6]. Their compression data and other single crystal results [7] suggested the presence of anisotropic fracture strength in NiA1 crystals. The (100) orientation was referred to as hard for a remarkably high yield strength and an elastic failure with no ductility. In contrast, some plastic deformation was observed in the soft (110) orientation. A similar conclusion was obtained by the tension tests recently [8-10]. The DBTT determined by the onset of plastic elongation Periodic

Table

for

B2 A J . m l n i d e s

27

26

Fe

!8

Co

Ni

1215C ~887A ' q - - - - - - L A ' ~

CO~'rANT

r6

r5

Re

Ru

Rh

Iloo¢ LO30A

LI)IIIIA

76

)7

Os

Pd

Ir

H0¢

Fig. 1. Transition metal elements forming B2 aluminides.

was found to depend upon the crystal orientation as well as the A1 concentration. The anisotropic character of plastic elongation to failure was attributable to the fact that the primary slip Burgers vector in NiA1 was (100). The resolved shear stress on dislocations becomes zero when the stress was applied along a (100) orientation. An elastic fracture definitely occurred without any plastic deformation as a result of the geometric constraint (Schmid factor). On the other hand, a higher fracture stress observed in (100) orientations might imply a better fracture resistance. Though stoichiometric FeAI crystals were relatively brittle, substoichiometric FeA1 intermetallics showed 2-3% plastic elongation in some crystalline orientations [11]. Our recent study on DS Fe-40A1 crystals indicated that a good Charpy impact energy was measured at low temperatures [12]. The ductility was believed to be associated with the (111) slip, which only occurred in those B2 compounds with a relatively low ordering energy. However, the elongation to failure at room temperature was limited to less than 5%, and fracture surfaces showed clear evidence of cleavage along specific crystalline planes. In this paper, the fracture properties of B2 aluminide single crystals including FeA1, NiA1 and CoAl are presented. The fracture behavior, especially the preferred cleavage plane as determined by using the fracture mechanics method, is described for each intermetallic compound. Possible explanations for the difference of cleavage planes in B2 aluminides are discussed. The anisotropy of fracture resistance in single crystals is attributable to the geometric constraint caused by the tendency of the crack to grow on the specific cleavage planes. EXPERIMENTAL

Materials

Three binary B2 aluminides that are constructed from the first row of transition metal elements in the periodic table are FeAI, CoAl and NiA1. Binary phase diagrams indicate that these B2 aluminides can form over a wide range of composition [13]. Alloys of nominal compositions (in at.%), Fe-40A1, Fe-48AI, Co-40A1, Co-49A1, Ni-46A1 and Ni-50A1 were selected for this study. A trace amount of Zr (0.02 at.%) was added as an impurity scavenger. Each composition was prepared by induction melting under argon employing high purity laboratory raw materials. The melts of 3 kg were cast into a chilled copper mold under a partial argon pressure. A master alloy of 1.2 kg cut from the ingot was used to grow a single crystal by the Bridgrnan process. The directional solidification unit consisting of a remelting furnace and a mold drawing furnace was modified to grow single crystals of 40 mm diameter by 150 mm long. A high purity alumina tube coupled with a grain selector was employed as the crystal mold. Laue X-ray diffraction results suggested that all aluminide crystals had a (100) growth direction.

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

Table 1. Room temperature fracture toughness of B2 alumm~de single crystals Fracture Alloy Crystal Notch toughness composition orientation direction IMPaw:'ml <001> 8.93 <1oo) Ni 50AI Cool> 80l <10o> Ni 50AI Ni 50AI <001> 7 29 <10o> <001) 8.62 <1oo> Ni-50AI 8.78 Ni 50A1 Ctoo> N1-50A1 <001> 9 31 CO0) (001) 7.41 (100> Ni-50AI <001) 7.95 <100> Ni-50AI <001> 4.53 <110> Ni 50AI Ni-50AI <001> 3 79 <110> 4.63 CII0> Ni-50AI (001> 470 CLIO> Ni 50AI ( -110> 441 (110> Ni 50AI Ni 50Al (-112> 5.12 C110> (-110) 5.14 Ni- 50A1 (-110> 491 <111) Ni-50Al COOl> 3.00 Cll0> Ni-46Al (001) 3 09 <11o> Ni 46A1 NiAI Alloy NA 12.87~ <1oo> NA 17 58" <1oo> NiAI Alloy NA 4.61" (11o> NiAI AHoy Co 49A1 ((1oi) 3 08 (11o) (001 > 2.20 Co 40A1 Fe-40AI C001) 33.3 <100) Fe-40A1 COOl) 33.4 <1oo> Fe-40AI <11o> COOl) 55.8 51.7 <110> Ve 40AI -I10) 375 <111> Fe 40AI 110> 43.4 <111> Fe 40AI 11.7 Directional solidified Ve 48A1 (near (100)) ~The result [29] is measured on specimens with a chevron notch instead of a through notch: the notch direction ~s not available (NA).

Fracture test Subsize bend bars with a rectangular cross section of 7.6 by 3.8 m m were used for fracture tests. An edge notch was cut by wire electro-discharge machining ( E D M ) at the center of the specimen. The notch plane was normal to the specimen axis. Figure 2 shows a schematic plot of the fracture specimen used in the study. All specimens were aligned by Laue X-ray diffraction so that both the specimen axis and the notch direction were along specified crystal orienrations. Attempts to precrack the notched specimens were not successful; the crack tended to deflect away from the notch plane in most cases. Therefore, notched specimens were employed and tested by 4-point bending, which would avoid the error associated with misalignment [14]. The loading rate was 0.0085 mm/s. Since all specimens failed elastically, the maximum load was used to calculate the fracture toughness, Kc, according to the equation A S T M E 399.

Fracture plane determination Cleavage occurred in every specimen tested in this study. Because each face of a specimen had a predetermined crystal orientation, the preferred fracture plane with low indices can be identified by examining different faces of the broken specimen. Flat and shiny facets on the fracture surface of broken specimens were also indexed directly by Laue X-ray diffraction. The specimen after fracture testing was mounted on a holder having two mutually perpendicular axes of rotation. The fracture surface was examined under an optical microscope. A high intensity light beam was aimed at the area of interest through an optical fiber tube mounted parallel to the optical axis of the microscope. The specimen was aligned by observing the reflection of the light beam. The fl.at facets were set in a horizontal position as the full reflection of the light beam was reached. The aligned specimen on the specimen holder was then moved to a Laue X-ray diffractometer for orientation determination. The accuracy of this optical alignment was within 2 ° .

P

P

A

,9 =

¢ P

k

RESULTS

Table 1 lists all r o o m temperature fracture toughness data measured in B2 aluminide crystals. The three B2 aluminides, FeA1, NiAI and CoAl, have room temperature fracture resistance significantly differing from one another. In a qualitative order, the fracture resistance of FeA1 > NiA1 > CoAl. All crystals fractured by cleavage at r o o m temperature, but the fracture behavior differs significantly from one another.

Fracture toughness Figure 3 shows the measured fracture toughness of Ni 50A1 single crystals along hard (100) and soft (110> orientations. The notch direction in both cases is (001). The temperature dependence of the fracture w toughness indicates the existence of a ductile-brittle transition. A substantial increase of Kc occurs when the temperature is raised above 200 C. Two important characteristics are noticed in the fracture of NiA1 crystals:

I[ o

I P

2729

=

Fig. 2. Four-point bend fraction specimen.

1. The fracture toughness in the hard (100) orientation is higher than that in the soft (110> orientation at all temperatures. The ratio is about 2 to 1 for all temperatures. The difference is explainable when the preferred cleavage plane is identified.

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

2730 30-

NiA1 SINGLE

CRYSTAL

FRACTURE TOUGHNESS 25,,.,.'"

L

20-

.x''"

[1oo] / /' / ....x............x" v 10, x..J ~ //~[11.0] U

o. 15.

o

5 , . -

; . . .

I00

:

. . .

200

: . .

300

•

:

-

400

.

.

:

500

TEMPERATURE, C

Fig. 3. Temperature dependence of fracture toughness measured in NiA1 single crystals. 2. The ductile-brittle transition as defined by fracture toughness occurs at the same temperature in both hard and soft orientations. Therefore, fracture toughness represents crystal's intrinsic resistance to cracking, one of the inherent crystal properties. In constrast, different DBTTs for different orientations have been observed for tensil ductility of NiA1 crystals [15]. Table 1 lists the measured toughness data for all B2 aluminide single crystals. The average fracture toughness measured in NiAI (111) crystals is 5.1 MPax/m, which is higher than in (110) but lower than in (100). The anisotropy of fracture toughness is considered to be associated with geometric constraints, the existence of preferred cleavage plane. A proposed model analogous to the Schmid factor of the resolved shear stress for single crystals will be discussed later. Also shown in Table 1, the substoichiometric Ni-46AI single crystals have a 30% lower fracture toughness compared with stoichiometric NiA1. The difference in the fracture toughness may be related to the defect hardening associated with the stoichiometry. NiAI intermetallics have a minimum room temperature hardness at the stoichiometric composition, Ni-50A1 [16]. Though a high strength crystal is expected to have a low fracture toughness, but limited data points do not allow the establishment of a valid strength/toughness relationship for NiA1 alloys. Considerable difficulty was experienced during machining of CoAl specimens because of surface cracking during mechanical grinding. Specimens were broken before reaching the final configuration for testing. In order to obtain meaningful values of fracture resistance, some crystals cut by wire-EDM were notched and tested without surface finishing. The recast layers on specimen surfaces were about

0.05 mm thick (0.25% of specimen thickness). The fracture toughness results measured at room temperature are also listed in Table 1. As expected, CoAl is much less tough than NiA1. The stoichiometric effect on the fracture toughness is observed again in CoAl single crystals; the near stoichiometric composition, Co-49A1 has a slightly higher fracture resistance than the Al-lean composition, Co-40A1. Attempts to grow single crystals of near stoichiometric Fe-48A1 were not successful. Fracture toughness was therefore measured on the directionally solidified crystal along the growth direction. The specimen failed elasticallry- and yielded a fracture toughness of 11.7 MPax/m. Unlike other two B2 aluminides, the room temperature hardness of FeA1 decreases as the A1 concentration is reduced from the stoichiometry. Therefore substoichiometric FeA1 crystals are expected to have a higher fracture toughness. The geometry of the specimen bars for Fe-40AI crystals was different from the others; the width was 5.1 mm instead of 7.6mm, and the thickness was kept the same (3.8mm). The results measured along three crystalline orientations: (100), (110) and ( I l l ) , are given in Table 1. The fracture toughness, Kc, was calculated from the maximum load in the test. In spite of the cleavage failure, Fe-40A1 single crystals exhibit much higher fracture toughness. However, these toughness data do not satisfy the ASTM minimum thickness requirement--the specimen thickness, B I> 2.5 (Kc/YS)2--for the valid plane strain fracture toughness; where YS is the yield strength. Indeed, the load-displacement curves of fracture toughness tests show a tearing tail rather than a catastrophic failure after passing the maximum load. A certain amount of plastic deformation occurs before fracture; therefore, the plain strain fracture toughness for Fe-40AI is expected to be lower than those reported in Table 1. The anisotropic fracture resistance in Fe-40AI is obvious since the three primary crystal orientations yield three different values of the fracture toughness. The minimum occurs in the (100) orientation in Fe-40A1 crystals, instead of in the (110) orientation as observed in NiA1 crystals. This distinction can be rationalized by the difference of the preferred cleavage plane as described below.

Cleavage plane Fracture test specimens contain an artificial defect (notch) to initiate the crack at a specific location. The crack path will deflect from the plane of maximum normal stress if there exists a preferred cleavage plane for the crystals. For bend specimens used in this study, the maximum normal stress occurs on the notch plane, which is perpendicular to the specimen axis. Therefore an ideal fracture test will occur when the normal of the cleavage plane coincides with the crystal specimen axis. The crack will grow on the notch plane with no deflection.

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

2731

Fig. 4. Fracture paths in NiA1 single crystals indicating a {110} cleavage plane. Figure 4 shows the sideviews of broken stoichiometric NiA1 specimens in (100) and (110) orientations. The ideal fracture test occurs in the (110) specimen. The fracture surface is flat and shiny optically. The river pattern on the cleavage surface was observed under a scanning electron microscope (SEM) as seen in Fig. 5. The step lines of the river pattern are relatively fine and discontinuous; some heavy step lines are observed with a large interspacing. On the other hand~ the (100) specimen shows a fracture path that deflects from the notch plane and follows the {110} surface trace as seen in Fig. 4, Instead of a sharp straight line, smooth curvatures are observed on the crack path. The fracture surface displays different areas along the crack path, and the curvature is associated with the transition from one

Fig. 5. River patterns on a NiAI {110} cleavage plane.

area to the other. Figure 6(a) shows that the crack initiated from the notch root tends to grow on an area consisting of stepwise layers before it moves to a {110} cleavage plane. These layers in the first part of the fracture path are termed the "transient" planes, forming a 45 ° angle with the crack growth direction projected on the {100} plane [Fig. 6(b)]. Two symmetric variants of the transient planes are observed. When the crack subsequently propagates on the {ll0} cleavage plane, only very fine river marks are observed on the flat fracture surface [Fig. 6(c)]. The fracture facets of the transient plane near the notch root appear to be crystalline planes because of their visual brightness. The surface trace analysis suggests that they are not planes with low indices. The optical alignment technique and Laue X-ray diffraction were applied to different areas on the fracture surface of (100) NiAl crystals. The results are plotted in a stereographic projection as shown in Fig. 7. Each layer structure consists of two sets of {51 l} fracture facets that are mirror symmetric about the (100) axis. These data are in agreement with the surface trace analysis, but the mechanism that forms fracture facets on the transient plane near the notch requires further investigation. The final cleavage on the {ll0} plane is confirmed again by the optical alignment technique. The ideal fracture test also occurs in the (110) specimens of high strength Ni-46A1 single crystals. The fracture surface that follows a {ll0} cleavage plane is flat without many features. Cleavage in CoAl crystals is seen when the specimens are broken during machining. Since the crystal

2732

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

r

Fig. 6. (a) Fracture surface of a NiAI ( 100)-oriented crystal, consisting of (b) fracture facets on {511} transient planes near the notch and (c) subsequently, the natural cleavage on

{ll0}.

was aligned in specific orientations, the cleavage plane observed on the fracture surface of unfinished specimens was identified as {110} planes. Therefore (110) specimens were prepared without surface grinding, and ideal fracture tests are performed in both Co-40A1 and Co-49A1 crystals. The fracture surface plane examined under SEM shows very few features as seen in Fig. 8. Figure 9 shows the fracture path that appeared on the Fe-40A1 crystals along three orientations: (100), (110) and (111). The notch directions are (001),

(001) and ( - 1 1 0 ) , respectively. The ideal fracture test occurs in the (100) specimen. Therefore, unlike other B2 aluminides, NiAI and CoAl, which cleaved on {110}, FeA1 has a {100} type cleavage plane similar to that of b.c.c, metals. A classical cleavage fracture with the river pattern is observed on the fracture surface as seen in Fig. 10(a). The step lines are well defined and tend to follow other variants of { 100} cleavage planes. The distinction of the features between {100} type and {110} type fracture surfaces is evident as revealed by the characteristics of the river pattern. In Fig. 9 the (111) Fe-40AI crystal has a fracture path tilted from the specimen axis. The surface trace analysis indicates that cleavage occurs on {100} as expected. In this case, the {100} cleavage takes place immediately from the notch root. The fracture facets on the {511} transient planes that are found in NiAI crystals are not observed in FeA1 crystals. Some secondary cracks formed along the primary crack in both (100) and (111) specimens, and their trace is confirmed to be other variants of {100} cleavage. The macroscopic fracture path in the (110) specimen seems to follow the notch plane, which is {110}. However, the crack formed in the (110) specimen is not so flat as in the other specimens. The SEM fractography, as seen in Fig. 10(c), reveals that the crack is composed of two variants of {100} cleavage. Alternative {100} fracture planes which develop a sawtooth crack front would create a macroscopic crack on {110} as observed. This fracture behavior is understandable when the geometrical constraints of this specimen are considered. For the ease of reference, let us define the notch plane is (110), and the notch direction is [001]; the crack front is then to be [ - 110]. Among three variants of {100} type cleavage plane, only the (001) plane contains the crack front. However, the resolved normal stress on the (001) plane is zero under the current configuration. Consequently, employing two other variants of cleavage planes, (100) and (010) can satisfy the constraint induced by the notch. A composite (110) fracture path contains the crack front [ - 110] and allows the crack to grow. This observation suggests that the defect nature, e.g. shape or orientation, in B2 aluminide crystals may affect the stress profile around it and therefore may alter the macroscopic crack path that is initiated from the defect. Nevertheless, the microscopic cleavage plane for intermetallic Fe-40AI single crystals is always {100}. It should be mentioned that the cleavage observed in B2 aluminides might be more complicated than the above results, which were obtained in the bend test with a simple geometry. The fracture specimen with different notch geometries and constraints have given different results than as described above. For example, bend fracture specimens with a chevron notch in (100) orientation do not show any evidence of {110} cleavage. The crack follows the path constrained by the notch, and the fracture surface shows

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

2733

NOTCH

DIRE[C(TI ON

oo~

o~o

.

.

.

.

o~o

~ A B~

off ~ 0 1 1 oo;

1100)

Fig. 7. Orientations of fracture facets observed in a NiAI (100)-oriented crystal (Fig. 6). only the fracture facets of {511 } transient planes. As a result the fracture toughness measured on chevronnotched specimens is higher than that on those with a through notch as seen in Table 1. However, the chevron-notched specimen in (110) orientation does show {110} cleavage and a similar fracture toughness as through-notched specimens. In case of smooth tensile specimens of NiAI crystals, the fracture surface not only shows {110} fracture facets, but also many different fracture facets with high indices, such as {511} transient planes. DISCUSSION Preferred cleavage plane Brittle fracture in single crystal intermetallics is anisotropic because of the existence of a preferred cleavage plane. Different cleavage planes have been observed in B2 aluminides, i.e. {110} for NiAI and CoAl, and {100} for FeAI, while the {511} transient planes have been observed in NiAI. The preferred cleavage plane for each B2 aluminide should be

200/~m

I

Fig. 8. River patterns on a Co-49A1 (110) cleavage plane.

predictable from the first principles of atomic bonding. Some relevant theories for the brittle fracture of a crystal are summarized briefly. When a perfectly brittle material fractures without any plastic deformation, Griffith's theory in its original form is applied. The fracture resistance is simply determined by the crack surface energy or by the free surface energy. A perfectly brittle single crystal is expected to cleave on the plane having a minimum surface energy. Theoretically, it is possible to calculate the surface energy of different crystalline planes provided that the atomic bonding energies are known. However, in a molecular dynamics simulation of the embedded atom method for NiA1 crystals, the prediction is not consistent with the experimental results. The {100} plane is predicted to have the lowest surface energy, while the {110} surface energy is about 60% higher [17]. The local atomic bond strengths in this simulation are determined by extended X-ray absorption edge fine structure spectra. Cottrell has proposed a dislocation mechanism for the nucleation of fracture in b.c.c, metals based on dislocation interactions [18]. The formation of sessile (100) dislocations occurs at the intersection of two {110} slip planes with ½1111) slip Burgers vectors. The opening of the crack in the {100} plane can be considered to result from the coalescence of a number of these (100) dislocations. A similar approach has been applied to FeA1, in which ½1111~ Burgers vectors become superdislocations in pairs [19]. This mechanism might be valid for the nucleation of a crack, but it is not capable of accounting for cleavage on the {100} plane as the crack propagates. In fracture toughness tests the notch can be set on any crystalline plane, but the cleavage does not follow the notch plane as seen in the above section. Gilman has proposed a simple criterion to explain why the b.c.c, transition metals cleave on {100}

2734

CHANG

[110]

et al.:

CLEAVAGE FRACTURE IN B2 ALUMINIDES

[111]

[100]

Fig. 9. Fracture paths in Fe~40A1 single crystals indicating a (100) cleavage plane. planes [3]. He considered a coordination number of 14 for the b.c.c, structure as recommended by Pauling instead of the more obvious 8. The crack surface energies of the {110} and {100} planes can be compared by counting the broken bond densities of both the first (8) and the second (6) nearest neighbors. Although there is a mistake in the calculation as pointed out in the Appendix, the idea of examining

~ - - ~O0~m

,,,1

the broken bonds on the cleavage plane offers a more comprehensive background for the future theoretical approach to this topic. In the B2 ordered structure each atom has 8 nearest neighbors of different atoms and 6 second nearest neighbors of similar atoms. Figure 11 compares the broken-bond structure of the {100} and {110} surfaces, On the { 100} cleavage plane, 4 nearest neighbor

,

2oo~m

,,,,,J

Fig. i0. Fracture surfaces of Fe-40AI single crystals oriented along (a) (100); (b) (111); (c) (110).

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

2735

is possible that they are related to the amount of plastic deformation prior to cleavage. Further investigation is being carried out. Toughness anisotropy

{ioo}

{11o}

Fig. 11. Crystalline planes {100} and {110} in a B2 lattice.

bonds and 1 second nearest neighbor bond are fractured for every atom. In constrast, only 2 nearest neighbor bonds and 2 second nearest neighbor bonds are broken when the cleavage occurs on the {110} plane. Different ratios of broken bonds may play a significant role in determining the preferred cleavage plane. The nature of bonds between the transition metal atom and the aluminum atom in B2 aluminides is the key parameter for physical and mechanical properties. A clear physical model of the electronic structure is not well defined at the present time. The NiA1 crystal has been the subject of most atomic bonding simulations [20-22]. In general, all theoretical calculations agree that Ni-A1 bonds are extremely strong and directional in NiAI crystals. There is a great discrepancy in theoretical models about whether Ni-AI bonds are ionic or covalent and if they are ionic, how the charge transforms. If we assume that the bonds are ionic and the bonding energy is proportional to the degree of ionization, a phenomenological explanation to the preferred cleavage plane in B2 aluminides is then possible. In those aluminides having a low ordering energy, such as FeA1, the fracture will occur on { 100} planes as in b.c.c, metals. As the ordering energy becomes stronger, i.e. the bonds become more ionic, {100} cleavage will be unfavorable since such a fracture creates two surfaces with different atoms, i.e. two charged surfaces with different signs. Surface charge will be balanced on the {110} cleavage planes. In addition, the atomically dense { 110} surfaces always have a minimum broken bond density. A similar empirical correlation has been made to explain the different Burgers vector observed in B2 intermetallics [23]. Consequently, the relationship between the cleavage plane and the Burgers vector exists: a {100} cleavage will occur for slip in a <111> direction, and a {110} cleavage for a <100> slip. The transition of slip directions from <111> to <100> has been reported in FeA1 at elevated temperatures or through alloying modification [24, 25]. Experimental verifications of the relationship between deformation and fracture characters are of great interest. No clear explanation for the occurrence of {511} transient planes in NiA1 can be given at this time. It

The brittle fracture of single crystals has been considered to be related to the resolved normal stress on the cleavage plane. Sohncke's law states that fracture occurs when the resolved normal stress reaches a critical value [26]. As verified by the fracture mechanics considerations in isotropic materials, fracture criteria depend not only on the stress but also on the size of the crack. Therefore, a fracture mechanics parameter, i.e. the stress intensity factor, instead of the fracture stress should be used. A revision of Sohncke's law is that fracture occurs when the resolved stress intensity factor reaches a critical value. This inherent crystal property can be measured by an ideal fracture test in which the cleavage plane is normal to the applied stress. For B2 aluminides, CoAl crystals have a value of 3 MPax/m; NiA1 crystals are 4.5 MPax/-m; and FeAI crystals are estimated above 12 M P a x / ~ . The anisotropy of fracture resistances in single crystals is simply a result of geometric constraints. The crack tends to follow a favorable cleavage plane, which may be oriented away from the direction of the applied stress and the orientation of the defect. Even if the stress is applied normal to the defect plane, the crack will not grow on the plane with a maximum normal stress unless it coincides with the cleavage plane. Fracture in other crystal orientations is expected to need some additional load higher than that for the ideal fracture test. The results of fracture tests in B2 aluminides verify the above concept of toughness anisotropy. The NiAI crystals that cleave on {110} planes have a minimum fracture resistance in the <110> oriented specimens. Similar minimum fracture resitance is measured from the <100> specimens of FeAI crystals that have a {100} cleavage plane. Vehoff recently reported a similar experimental observation of anisotropic fracture resistances in NiAI single crystals [27]. The {110} cleavage plane was observed in both <100) and <110> specimens. Vehoff compared his experimental data with the theoretical calculation of the cleavage energy made by Yoo and Fu using a first principle total energy approach [28]. The correspondence between the measured and predicted values for the fracture toughness of NiA1 <110> crystals is remarkably good. However, the theoretical value for FeA1 seems too low compared with the experimental results reported above. Furthermore the theoretic calculation predicted a {110} cleavage plane for both FeAI and NiAI. This does not agree with our experimental observation that Fe-40AI has a { 100} cleavage plane. The difference of {100} and {110} cleavage energies for FeA1 in Yoo's calculation is relatively small. Incorporating the crack tip plasticity into the fracture

2736

CHANG et al.: CLEAVAGE FRACTURE IN B2 ALUMINIDES

energy calculation will be essential for developing a satisfactory theoretical model; stoichiometry difference may also be playing a role here. Before the precise formula is available for the calculation of the stress intensity factor in anisotropic single crystals, an approximation can be made by calculating the resolved normal stress intensity on the cleavage plane. The fracture resistance in any given orientation is correlated with the fracture toughness measured under the ideal fracture test by a normal ratio r = cos 2 ~b; where ~b is the angle between the normal of the cleavage plane and the axis of applied stress. The normal ratio for NiA1 crystals that cleave on {110} planes is plotted over the entire standard stereographic triangle in Fig. 12. The (100) specimen is predicted to have the highest fracture resistance, and the normal ratio is r = 0.5. Experimental data of NiA1 single crystals shown in Fig. 3 support such an oversimplified approximation. The measured fracture resistance in ( I 0 0 ) specimens is about twice of that measured in (110) specimens at all temperatures. Additional works on the fracture toughness of other crystal orientations are in progress, and the preliminary results indicate a good agreement with the approach using the criterion of critical resolved normal stress intensity (CRNK). Figure 13 shows some results measured on specimens cut from the same single crystal; each orientation has duplicate tests. The calculated values represented by the circles are obtained by multiplying C R N K with the normal

~

,2,

R

j

10

d

v

a r.=,,,

J ]

, ~

I

I

n T®,t#2 -o- Normal Kc

i I •

// // /z

::

//: .//1

i// o" [loo] 0rienfafion Fig. 13. Anisotropy of fracture toughness in NiAI single crystals. ratio (r) for each orientation. The fit of experimental and theoretical data is reasonably good. The occurrence of the {511} transient plane near the notch in NiA1 crystals might further complicate the situation of fracture anisotropy. The mechanism of forming these transient planes is not available yet, and its influence on the measured fracture resistance is currently being studied. The stoichiometric effect on the fracture toughness of B2 aluminides can be explained by the correlation of strength and toughness. A higher fracture toughness is measured in those whose strength is lower. Both NiA1 and CoAl has a minimum hardness at the stoichiometric composition, and therefore a maximum fracture toughness is obtained. In contrast, Al-lean FeA1 is weaker than stoichiometric FeA1; a higher fracture toughness is measured when the AI concentration is reduced. CONCLUSIONS

111o.667 /'~

WITH A {110~ CLEAVAGE P L A N E / ~

2y,

100

NiAI Single Crystal

: COS2~

FRACTURE NORMALIZING RATIO FOR A CUBIC CRYSTAL

o.s~

14

/Io.,

I

I°.8

I

/

1

I

310

210

1

/0.975 , / 520

=,.o 110

Fig. 12. Contours of constant fracture normalizing ratio in the standard stereographic triangle for {i10} cleavage.

1. B2 aluminides, including FeAI, CoAl and NiAI, fracture by cleavage on a preferred crystalline plane at room temperature. Both NiA1 and CoAl single crystals cleave on {110} planes, while FeAI cleaves on {100} planes. For NiAI, fracture on {511} transient planes may proceed before cleavage on {110} planes. 2. Anisotropy of the fracture resistance in B2 aluminide crystals is attributable to the presence of preferred cleavage planes. A high fracture resistance is measured in the orientation away from the cleavage plane. The critical resolved normal stress intensity (CRNK) of a B2 aluminide crystal is obtained by an ideal fracture test. CoAl is 3 MPa,q/m; NiA1 is 4.5 MPax/~; and FeAl is about 12 MPax/~. 3. The fracture toughness of in NiA1 crystals starts to increase substantia!]y as the temperature is increased above 200°C, independent of the crystal orientation. 4. A stoichiometric effect on the fracture toughness has been observed in all B2 aluminides studied. The measured fracture toughness is related inversely with the alloy strength.

C H A N G et al.: CLEAVAGE F R A C T U R E IN B2 A L U M I N I D E S Acknowledgements--The authors would like to thank R. Smith, R. A. Rosa, D. A. Catharine, and D. W. Marsh for their technical assistance with the experiments. Useful discussions with R. D. Field, D. Lahrman and J. E. Hack (Yale University) were greatly appreciated.

REFERENCES I. N. P. Allen, B. E. Hopkins and J. E. McLennan, Proc. R. Soc. A 234, 221 (1956). 2. C. F. Tipper, The Brittle Fracture Story. Cambridge Univ. Press, New York (1962). 3. J. J. Gilman, Fracture of Solid, p. 541. Interscience, New York (1962). 4. I. Baker and P. R. Munroe, High Temperature Aluminides and Intermetallics (edited by S. H. Wang et al.), p. 425. TMS, Warrendale, Pa (1990). 5. M. H. Yoo, T. Takasugi, S. Hanada and O. Izumi, Mater. Trans. Japan Inst. Metals 31, 435 (1990). 6. R. T. Pascoe and C. W. A. Newey, J. Metal Sci. 2, 138 (1968). 7. A. Ball and R. E. Smallman, Acta metall. 14, 1349 (1966). 8. R. D. Field, D. F. Lahrman and R. Darolia, Mater Res. Soc. Symp. Proc. 213, 255 (1991). 9. D. F. Lahrman, R. D. Field and R. Darolia, Mater Res. Soc. Symp. Proc. 213, 603 (1991). 10. R. D. Field, D. F. Lahrman and R. Darolia, Acta metall, mater. 39, 2951 (1991). 11. Y. Umakoshi and M. Yamaguchi, Phil. Mag. A 41,573 (1980). 12. K.-M. Chang, Metall. Trans. 21A, 3027 (1990). 13. T. B. Massalski, Binary Alloy Phase Diagrams. Am. Soc. Metals, Metal Park, Ohio (1986). t4. K.-M. Chang and D. A. Catharine, Rapid Fracture Toughness Measurements, G E - C R D Report 90CRD 154 (1990). 15. R. Darolia, JOM 43-3, 44 (1991). 16. J. H. Westbrook, Metall. Rev. 9, 145 (1965). 17. P. C. Clapp, M. J. Rubins, S. Charpenay, J. A. Rifkin, Z. Z. Yu and A. F. Voter, Mater. Res. Soc. Symp. Proc. 133, 29 (1989). 18. A. H. Cottrell, Trans. Am. Inst. Metall. Engrs 212, 192 (1958). 19. P. R. Munroe and I. Baker, Acta metall, mater. 39, 1011 (1991). 20. S.-C. Lui, J. W. Davenport, E. W. Plummet, D. M. Zehner and G. W. Fernando, Phys. Rev. B. 42, 1582 (1990). 21. R. Darolia, R. D. Field and D. F. Lahrman, Alloy Modeling and Experimental Correlation for Ductility Enhancement in Near Stoichiometrie Single Crystal Nickel Aluminide, A F O S R / N E Report F49620-88-C0052 (1990).

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22. A. G. Fox and M. A. Tabbernor, Acta metall, mater. 39, 669 (1991). 23. W. A. Rachinger and A. H. Cottrell, Acta metall. 4, 109 (1956). 24. D. K. Partick, K.-M. Chang, D. B. Miracle, and H. A. Lipsitt, Mater. Res. Sac. Syrup. Proc. 213, 267 (1991l. 25. M. G. Mendiratta, H. Kim and H. A. Lipsitt, Mctall. Trans. 15A, 395 (1984). 26. A. Deruyttere and G. B. Greenough, J. Inst. Metals 84, 337 (1955 56). 27. H. Vehoff, Proc. NATO Adranced Research 14')~rkshop, Ordered Intermetallics--Physical Metallurgy and Mechanita/Behavior, June 23-29, Irsee, FRG, in press. 28. M. H. Yoo and C. L. Fu, Scripta metall, mater., in press. 29. K. Bain and R. Darolia, unpublished work, GE Aircraft Engines, Evendale, Ohio.

APPENDIX Broken-Bond Densities in b.c.c. Lattices The following treatments can be applied to either b.c.c, or B2 structures. Each atom in the lattice has 8 same nearest neighbors at a distance of 0.866 a 0 and 6 same second nearest neighbors at a distance of a0; where a 0 is the lattice parameter. Let A and B be the energies of broken nearest and second nearest neighbor bonds, respectively. (In the ordered B2 structure A represents the bonding between unlike atoms, and B represents the averaged bonding between like atoms.) The broken-bond density of a crystalline plane is defined as the energy per unit area is required for cleavage on this plane. Referring to Fig. 11, the broken-bond density of the {I00} plane is D~0~,= (4A + 2B)la~

(AI)

where the atom on the first layer on a ~100} fracture plane contributes broken bonds of 4A and I B, and the atom on the second layer contributes lB. On the {1101 plane, only the atom on the first layer contributes broken bonds of 2.,t and 2B, and the broken-bond density is DIto -- ~'v' _(A + B)/a~.

(A2~

The difference of the fracture energies between 't 100~ and { 110} planes is Dr,,0 - Dt~0 = 2(x,/-2 - 1)(, .' 2,4 - B),,t~.

(A3)

The above value is always greater than zero since A > B in either b.c.c, or B2 structures. Therefore the dense packed [110} plane has a low broken-bond density. (Gilman used 4A + B instead of4A + 2B in equation (AI), which rcsultcd in a different conclusion [3].)