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Nanometer-scale crack initiation and propagation behavior of Fe3Al-based intermetallic alloy

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Nanometer-scale crack initiation and propagation behavior of Fe3Al-based intermetallic alloy
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  Nanometer-Scale Crack Initiation and Propagation Behavior of Fe3AI-Based Intermetallic Alloy L.J. QIAO, X. MAO, and C.Z. CHEN The initiation and propagation of nanometer-scale cracks have been investigated in detail by in situ transmission electron microscope (TEM) observations for the intermetallic compound Fe3A1 under mode I loading. No dislocation was detected and no dislocation emission was found when cracks propagated directly from the thin edge of a double-jet hole where the thickness of the foil was below a critical thinness. Thinning took place in the thicker region of the foils because a great number of dislocations were emitted from the crack tip, and then an electron semitransparent region was formed in front of the crack tip. Following this process, a dislocation- free zone (DFZ) was formed. The maximum normal stress occurs in the zone. Nanometer-scale cracks initiated discontinuously ahead of the main crack tip in the highly stressed zone. The size of the smallest nanocrack observed was about 3 nm, and the tip radius of the nanocracks was less than 1 nm when the applied loading was low. The radius of the main crack tip was about 2.5 nm. The distances between discontinuous nanocracks and the main crack tip were about 5 to 60 nm, depending on the applied tensile loading. A relationship was found between the tensile loading and the nanocrack distance from the crack tip. The distance increases with the tensile loading, which is consistent with an "elastic-plastic" theoretical model. I. INTRODUCTION FRACTURE involves the breaking of atomic bonds during crack propagation. The high local stresses at some microregions should, therefore, reach the theoretical level equivalent to atomic cohesion during the fracture. Usu- ally, the applied stress is much lower than the theoretical stress. The pileup of dislocations would result in a stress concentration that produces stresses sufficient for atomic decohesion, in'31 A few models have been suggested based on combined studies of fracture mechanics and discre- tized dislocation microstructural analysis, t4-Tj The idea of dislocation shielding was introduced by Rice and Thomson.t6j They concluded that crack propagation would be by means of cleavage if dislocation emission was dif- ficult; otherwise, the propagation would be ductile. A number of studies have been made in an effort to under- stand the mechanisms controlling ductile and brittle frac- ture. t7-131 Many of these studies have focused on dislocation emission from the crack tip and the concept of a dislocation-free zone (DFZ). Elastic-plastic models have been established to calculate the stress level and the stress distribution in front of the crack tip by com- puter simulationJ 4,5,~4-J6~ The stress contributions have been calculated for cases with and without dislocation emission by Lii et al. [4] Huang and Gerberich, t~ Marsh et al. t~Sj and Chen et al. llrl The position of maximum normal stress will be located at the crack tip if dislo- cation is not emitted. The situation can be described by linear elastic theory. This position then shifts away from L.J. QIAO, Postdoctoral Research Associate, and X. MAO, Associate Professor, are with the Department of Mechanical Engineeging, The University of Calgary, Calgary, AB T2N 1N4, Canada. C.Z. CHEN, Lecturer, is with the Department of Materials Physics, University of Science and Technology-Beijing, Beijing 100083, People's Republic of China. Manuscript submitted June 6, 1994. the crack tip due to dislocation emission. The shifted distance increases with the increase of the applied stress- intensity factor Kl.t4'51 There are a few instances, however, of direct experi- mental observations of crack initiation. The transmission electron microscope (TEM) in situ tensile observations have been performed on austenitic stainless steel and the intermetallic compounds Ti3AINb and TiA1. tZ7'lsA91 The results show that dislocations were emitted from the crack tip followed by the formation of a highly stressed zone. The average strains of the region ahead of the crack tip were about 0.1, 0.07, and 0.02 for 310 stainless steel, Ti3A1Nb, and TiAl, respectively, t17,18,19] At such strains, a crack with a nanometer-scale size initiated preferen- tially in the highly stressed zone. The nanometer crack would blunt into a void for 310 stainless steel, which resulted in ductile fracture. For Ti3AINb and TiA1, how- ever, the nanometer crack would propagate continuously and join up with the main crack. For these two inter- metallics, this resulted in brittle fracture. Fe3A1 intermetallic compounds developed in recent years have a large potential for high-temperature appli- cations. However, the alloy has limitations in such ap- plications due to brittle behavior at room temperature. A number of efforts have been made to understand the nature of the fracture behavior in the intermetallic alloys. F2~ The macroscopic fracture surfaces are gen- erally formed by cleavage, although the alloy exhibits significant ductility (-15 pct) in air. t2~ The crack nu- cleation and propagation are still unclear for the inter- metallics. It is necessary to study the fracture process, which includes crack initiation and propagation. The ob- jectives in this study are as follows: (1) to observe nano- crack initiation and propagation of the Fe3A1 alloy in the in situ TEM tensile condition and (2) to study the mech- anism of the nanocrack initiation and propagation for the intermetallics. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JUNE 1995--1461  II. EXPERIMENTAL PROCEDURES The Fe3A1 alloy used in this study was supplied by Oak Ridge National Laboratory (Oak Ridge, TN). The compositions are shown in Table I. Preparation was by vacuum induction melting, extrusion to 25-ram thick- ness, and rolling to a final thickness of 6.4 mm. The alloy was heat-treated at 700 ~ for 1 hour, followed by air cooling. The heat treatment produced a mainly B2 ordered structure accompanied by some amount of DO3 order. Sheets 0.2-mm thick were cut from the bulk ma- terial. These sheets were mechanically thinned to about 60/~m using emery paper and then electropolished by a double-jet apparatus in a methanol solution containing 10 pct H2SO4 to obtain thin-foil specimens for the TEM. Small cracks would appear around the double-jet hole due to the preparation. Some hydrogen is probably in- troduced during the process, and this will be discussed in a separate article. 1251 A Hitachi H800 TEM was used in the in situ study at an applied acceleration voltage of 175 kV and a tensile loading rate of 50 nm/s. The ap- plied loading is controlled by a certain displacement through tensile time. Step tension was applied during in situ tensile observation. The displacement increment was about 5 nm for one step. Stresses around the double- jet hole and the cracks cannot be calculated exactly be- cause of the complicated shape of the double-jet hole and the random array of the small cracks. The applied stress was established qualitatively. the thinner regions that were below a critical thinness, e.g., at some edge of the double-jet hole, is character- ized by the absence of a plastic zone ahead of the crack tip. That is, no dislocation pileups or dislocation tangles were detected during tension in the region immediately adjacent to the edge of the crack. Only elastic strain con- tours appear around the crack tip, as shown in Figure 3. This type of propagation appears to be brittle, although a cleavage plane could not be identified. Also, grain boundaries are not the weakest interface. They do not fracture even if the interface is perpendicular to the ten- sile direction, as indicated by the arrows labeled "GB" in Figure 3. In the thicker regions, a great number of dislocations were emitted from the crack tip when ten- sion was applied. Thinning occurred ahead of the crack tip, which resulted in a semitransparent region. The local stress intensity for dislocation emission Kie increases, and the dislocation emission becomes more difficult as dis- locations pile up in front of the crack tip. A highly stressed III. EXPERIMENTAL RESULTS A. General Observation In an air environment, strain to fracture was found to be 14.3 when tested with a strain rate over 2 x 10-4/S. It was 10.4 when the strain rate was reduced to 1 x 10-6/s. However, the fracture surfaces are mainly cleav- age surfaces accompanied by some intergranular zone at both fast and slow strain rates, as shown in Figure 1. The fracture surfaces are brittle, though the alloy exhib- ited significant ductility. The interrnetallic compound Fe3A1 used in this inves- tigation is of the ordered B2 structure accompanied by some amount of DO3 order. An ordered diffraction pat- tern is obtained by the selected area diffraction (SAD) method, as shown in Figure 2 ({110} diffraction pattern). The alternate bright and dim spots are the distinguishing features of the ordered state. These dim spots will dis- appear in the disordered state. There are many super- dislocations and subgrain boundaries in the alloy. These play an important role in the process of crack initiation and propagation. The fractures were of two types: continuous and dis- continuous. Here, continuous means that a nanocrack initiates from the main crack tip, while discontinuous implies initiation ahead of the tip. Crack propagation in Fig. 1 --Transgranular fracture surface containing some intergranular zone in air with strain rate of 2 x 10-3/s. Table I. The Compositions of Fe3AI Based Alloy Elements Cr Nb Mo B C Zr A1 Fe At. Pct 5.0 0.5 0.5 0.2 0.2 0.1 28.0 balance Fig. 2--{110} zone axis SAD pattern of the Fe3AI alloy with the B2 ordered structure accompanied by DO3 order. 1462--VOLUME 26A, JUNE 1995 METALLURGICAL AND MATERIALS TRANSACTIONS A  Fig. 3--Crack propagation from the thin edge of a double-jet hole and strain contours. No dislocations were detected. GB--grain boundary; CT--crack tip. (a) A crack initiated from the edge. (b) The crack appeared to propagate directly without dislocation emission. zone was formed in the semitransparent region when equilibrium was achieved. Following this, a nanometer- scale crack initiated in the zone in which the equilibrium ceased upon nanocrack initiation. Dislocations were emitted again, which resulted in the blunting of the nanocrack. Crack propagation occurred by the initiation and the growth of nanometer cracks ahead of the main crack tip in the electron semitransparent region. With the sizes of nanocracks about 3 to 70 nm, the internanocrack distances are about 5 to 100 nm. It was found that the size and the distance depended on the level of the ap- plied stress-intensity factor Kt. B. In Situ Observations of the Discontinuous Nanocrack Initiation The position of maximum stress will be located at the crack tip if dislocation emission does not occur when a tensile stress is applied. As a result, the crack propagates directly from the crack tip (Figure 3). However, the crack tip is no longer the position of maximum stress if dis- locations were emitted from the crack tip. This is due to the shielding stress from dislocations. This causes the position of maximum stress to shift away from the main crack tip. Ia,51 The shift distance is dependent on the level of the applied K~. S~ The theoretical results in elastic-plastic models indicated that the maximum stress would ap- proach the theoretical atomic bond strength 14-6,261 when the applied K, is over a critical value. Selected area dif- fraction results showed that the elastic strains in the highly stressed zone ahead of the crack tip were 0.1 ,[LTI 0.07,t~8~ and 0.02 ~191 for 310 stainless steel, Ti3AINb, and TiA1, respectively. They approached the theoretical fracture stress o , which is about E/10, I27] where E is Young's modulus. Therefore, nanocrack initiation, growth, and coalescence are possible in the region ahead of the main crack tip. Figure 4 illustrates this dynamic process of the initi- ation, growth, and coalescence of the nanocrack. Figure 4(a) shows a crack tip A (marked by an arrow) with a radius of 2.5 nm. The point C is a site of dis- continuous nanocrack initiation. The foil specimen is semibroken in the region AB in front of the main crack tip. The crack does not penetrate the foil in the direction perpendicular to the foil plane. This shows that the stress distribution is not uniform on the crack plane ahead of the crack tip, though it is in the thinned region. As shown in Figure 4(b), the nanocrack C (-3 nm) grew and con- nected with the main crack A. Another discontinuous nanocrack initiated at the point D. The size of nanocrack D is about 3 nm. This is the smallest crack reported as yet in the literature. The distance between C and D is about 5 nm. The main crack blunted and opened ac- cordingly as the nanocrack D grew, and then the cracks coalesced, as shown in Figure 4(c). The crack tip is very sharp at the point of coalescence, with a crack tip radius of less than 1 run. Then the crack blunted with time under constant displacement, as shown in Figure 4(d). Suc- cessive propagation consisted of the initiation, growth, and coalescence of nanocracks. The detailed process of nanocrack initiation in the highly stressed zone is shown in Figure 5. A small bright spot formed ahead of the crack labeled "B" in Figure 5(a). It is difficult to distinguish the spot because of its slight contrast with the bulk. It became larger and brighter with time, as shown in Figure 5(b). It can be seen in Figure 5(c) that the bright spot has developed into a nanocrack. The nanocrack grew and connected with the propagating main crack, as shown in Figure 5(d). C. Effect of Tensile Loading on the Position of the Nanocrack Initiation The loading is low in the observations discussed in Section B. The initiation, growth, and coalescence of nanocracks were relatively slow under low loading. The process could be observed and recorded in detail. Dis- continuous initiation behavior was different when the ap- plied loading was high. The main differences between METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JUNE 1995-- 1463  Fig. 4-- Successive stages of discontinuous nanocrack initiation, growth, and coalescence with main crack. The arrows A, B, C, and D indicate fixed position in the foil. (a) Blunted crack A, partial fracture region AB, and discontinuous nanocrack C; (b) the connection between nanocrack C and main crack A, and the initiation of new discontinuous nanocrack D with a size of 3 nm, the distance between "C" and "D" is 5 nm; (c) the blunting of crack tip C and the connection between C and D; and (d) blunted crack tip D. high and low loading are the position of initiation and the size of discontinuous nanocracks. The distances from the main crack to nanocracks and the nanocrack size in- creased as the loading was increased, as shown in Figure 6. In Figure 6(a), three nanocrack positions that will form are indicated. The nanocrack A formed at the low loading just connected with the main crack, and the cracks B, C, and D initiated as the loading was in- creased. The nanocrack B connected with the main crack, and the points C and D became thinner and brighter, as in Figure 6(b), but did not yet penetrate through the thickness. They then formed in Figure 6(c). The nano- crack D with a length of 25 nm initiated and propagated along a cleavage plane. Another nanocrack, E, formed in Figure 6(d). The distances from the main crack tip to the center of the nanocracks increased with increased loading, as shown in Figure 6(d). The distance increased from 5 to 100 nm. This showed that the size of the nanocrack at high loading is larger than that at low loading. Similar results were found for in situ TEM observations of high-purity gold foils. ~2s] The results are consistent with the theo- retical model of Gerberich and co-workers. 14,29] In this model, the position of maximum stress was shifted away from the main crack tip due to the shielding stress of dislocations emitted in front of the crack tip. The higher the applied stress-intensity factor Kt, the farther the po- sition of maximum stress shifted away from the crack tip. IV. DISCUSSION A. The Continuous Propagation from the Main Crack Tip The maximum stress will be located at a crack tip for an elastic solid with the crack propagating from the tip under tensile loading. In Figure 3, the crack propagated continuously from the thin edge of the double-jet hole. No dislocations were detected, and no dislocation emis- sion was found near the crack tip during the process. Only strain contours appeared around the crack tip. This process belongs to elastic fracture, which is perhaps due to the small film thinness around the double-jet hole in the foils. The shear stress for multiplication of dis- location is related to the thinness of the foil. It has been 1464--VOLUME 26A, JUNE 1995 METALLURGICAL AND MATERIALS TRANSACTIONS A  Fig. 5--The detailed successive process of nanocrack initiation ahead of a main crack tip. (a) A nanocrack will initiate at point B. (b) Point B is bigger and brighter. (c) Nanocrack formation is shown here. (d) Connection with the main crack occurs here. reported that cracks propagate without the involvement of dislocations if the thinness of the foils is below a crit- ical value) 27J The critical shear stress for the operation of a Frank-Read source can be given by ~'FR = 2/xb/l t3~ in the thin region with a thinness l. Here /x and b are shear modulus and Burgers vector, respectively. If a maximum working shear stress rmax --< 2txb/I dis- location emission would not take place in the thin region l. At a given maximum shear stress, the thinness l re- quired for dislocation emission is given by 2~b l -~- [11 "Tma Brittle fracture will occur without dislocation emission if the normal stress Cryy eaches the theoretical fracture stress o% before the maximum working shear stress is over the critical shear stress ~'FR. The theoretical fracture stress is about E/10.1271 For the Fe3A1 alloy, cr, h is about 14,100 MPa. 12~ The stress distribution in front of the crack can be cal- culated from elastic fracture mechanics for the case of Figure 3. As can be found from Figure 3, the crack ra- dius p is about 3 nm. Provided that the root radius of the crack is small, the stresses near a conical crack tip are similar in form to those of the crack with p = 0, but the srcin of r is located at a distance p/2 outward from the center of the crack root, radius p, as shown in Figure 7. The stresses in front of the crack tip are given by Eqs. [2] through [5]. [311 [ os (1 s n~ sin P cos-~] [2] 2r ~,,y - cos 1 + sin- sin 2 + -- cos [3] 2r % - sin - cos - cos sin [4] ~ 2 2 2 2r Tmax = -]- O'xx- O yy [51 METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JUNE 1995-- 1465
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