AnIn Situ transmission electron microscope deformation study of the slip transfer mechanisms in metals

The slip transfer mechanisms across grain boundaries in 310 stainless steel, high-purity aluminum, and a Ni-S alloy have been studied by using thein situ transmission electron microscope (TEM) deformation technique. Several interactions between mobile lattice dislocations and grain boundaries have been observed, including the transfer and generation of dislocations at grain boundaries and the nucleation and propagation of a grain boundary crack. Quantitative conditions have been established to correctly predict the slip transfer mechanism. This paper is based on a presentation made in the symposium “Interface Science and Engineering” presented during the 1988 World Materials Congress and the TMS Fall Meeting, Chicago, IL, September 26–29, 1988, under the auspices of the ASM-MSD Surfaces and Interfaces Committee and the TMS Electronic Device Materials Committee.     

Neutron and X-ray Microbeam Diffraction Studies around a Fatigue-Crack Tip after Overload

An in-situ neutron diffraction technique was used to investigate the lattice-strain distributions and plastic deformation around a crack tip after overload. The lattice-strain profiles around a crack tip were measured as a function of the applied load during the tensile loading cycles after overload. Dislocation densities calculated from the diffraction peak broadening were presented as a function of the distance from the crack tip. Furthermore, the crystallographic orientation variations were examined near a crack tip using polychromatic X-ray microdiffraction combined with differential aperture microscopy. Crystallographic tilts are considerably observed beneath the surface around a crack tip, and these are consistent with the high dislocation densities near the crack tip measured by neutron peak broadening. This article is based on a presentation given in the symposium entitled “Neutron and X-Ray Studies for Probing Materials Behavior,” which occurred during the TMS Spring Meeting in New Orleans, LA, March 9–13, 2008, under the auspices of the National Science Foundation, TMS, the TMS Structural Materials Division, and the TMS Advanced Characterization, Testing, and Simulation Committee.     

A dislocation shielding prediction of the toughness transition during cleavage of semibrittle crystals

An interconnected set of observations assesses current equilibrium models of the ductile-brittle-transition temperature (DBTT). This involvesin situ transmission electron microscopy (TEM) studies of crack-tip dislocations in single and polycrystals and bulk fracture toughness tests at various temperatures. Beyond K_I values of 8 MPa · m^1/2 in both iron-base single and polycrystals, large numbers of redundant dislocations are created, as postulated recently by Weertman. [38] Still, the necessary shielding dislocations, as required by equilibrium, can be detected at values as high as 20 and 40 MPa · m^1/2 byex situ TEM and electron channeling, respectively. In addition, the close approach of dislocations to the crack tip in some of the studies, as opposed to others, suggests that large dislocation free zones (DFZ) are a thin-film artifact. However, a failure criterion based partly on the Rice-Thomson model’21 is both consistent with the absence of a large DFZ and observed fracture toughness variations with test temperature. It is emphasized that this toughness transition is entirely in the semibrittle regime where cleavage is the failure mode. Nevertheless,K _lc values increase from 3 to 60 MPa·m^1/2 with an increase in test temperature. This article is based on a presentation made in the symposium “Quasi-Brittle Fracture” presented during the TMS fall meeting, Cincinnati, OH, October 21–24, 1991, under the auspices of the TMS Mechanical Metallurgy Committee and the ASM/MSD Flow and Fracture Committee.     

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

The initiation and propagation of nanometer-scale cracks have been investigated in detail byin situ transmission electron microscope (TEM) observations for the intermetallic compound Fe₃Al 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.     

Dislocation structure behind a shock front in fcc perfect crystals: Atomistic simulation results

Large-scale molecular dynamics simulations are used to investigate the dislocation structure behind a shock front in perfect fcc crystals. Shock compression in both the 〈100〉 and 〈111〉 directions induces dislocation loop formation via a sequential emission of partial dislocations, but in the 〈100〉 case, this process is arrested after the first partial, resulting in stacking-fault loops. The large mobility of the bounding partial dislocations results in a plastic wave that is always overdriven in the 〈100〉 direction; the leading edges of the partials are traveling with the plastic front, as in the models of Smith and Hornbogen. In contrast, both partials are emitted in 〈111〉 shock compression, resulting in perfect dislocation loops bounded only by thin stacking fault ribbons due to the split partial dislocations. These loops grow more slowly than the plastic shock velocity, so new loops are periodically nucleated at the plastic front, as suggested by Meyers. This article is based on a presentation given in the symposium “Dynamic Deformation: Constitutive Modeling, Grain Size, and Other Effects: In Honor of Prof. Ronald W. Armstrong,” March 2–6, 2003, at the 2003 TMS/ASM Annual Meeting, San Diego, California, under the auspices of the TMS/ASM Joint Mechanical Behavior of Materials Committee.     

Work-hardening and recovery mechanisms in gamma-based titanium aluminides

The work-hardening mechanisms in two-phase γ-titanium aluminide alloys were characterized in terms of the glide obstacles determining the velocity and slip path of dislocations, utilizing transmission electron microscopy (TEM) observations and thermodynamic-glide parameters. There was clear evidence that short-range obstacles in the form of dislocation debris and dipoles were produced during plastic deformation at room temperature. These dislocation obstacles contributed significantly to work hardening. The observed strong strain hardening arose from long-range elastic dislocation interactions and the production of dipole and debris defects. The thermal stability of these deformation-induced defects was assessed by isothermal and isochronal annealing. The results indicated that the dipole and debris defects were relatively unstable upon annealing at moderately high temperatures, which led to significant recovery of work hardening. This article is based on a presentation made in the symposium entitled “Fundamentals of Structural Intermetallics,” presented at the 2002 TMS Annual Meeting, February 21–27, 2002, in Seattle, Washington, under the auspices of the ASM and TMS Joint Committee on Mechanical Behavior of Materials.     

Shock Compression of Monocrystalline Copper: Atomistic Simulations

Molecular dynamics (MD) simulations were used to model the effects of shock compression on [001] and [221] monocrystals. We obtained the Hugoniot for both directions, and analyzed the formation of a two-wave structure for the [221] monocrystal. We also analyzed the dislocation structure induced by the shock compression along these two crystal orientations. The topology of this structure compares extremely well with that observed in recent transmission electron microscopy (TEM) studies of shock-induced plasticity in samples recovered from flyer plate and laser shock experiments. However, the density of stacking faults in our simulations is 10² to 10⁴ times larger than in the experimental observations of recovered samples. The difference between experimentally observed TEM and calculated MD results is attributed to two effects: (1) the annihilation of dislocations during post-shock relaxation (including unloading) and recovery processes and (2) a much shorter stress rise time at the front in MD (<1 ps) in comparison with flyer-plate shock compression (∼1 ns). This article is based on a presentation made in the symposium entitled “Dynamic Behavior of Materials,” which occurred during the TMS Annual Meeting and Exhibition, February 25–March 1, 2007 in Orlando, Florida, under the auspices of The Minerals, Metals and Materials Society, TMS Structural Materials Division, and TMS/ASM Mechanical Behavior of Materials Committee.     

Dislocation simulation of brittle-ductile transition in ferritic steels

Two-dimensional discrete dislocation simulations of the crack-tip plasticity of a macrocrack-micro-crack system representing the fracture behavior in ferritic steels are presented. The crack-tip plastic zones are represented as arrays of discrete dislocations emitted from crack-tip sources and equilibrated against the friction stress. The dislocation arrays modify the elastic field of the crack; the resulting field describes the elastoplastic crack field. The simulated crack system involves a microcrack in the plastic zone of the macrocrack (elastoplastic stress field). The effects of the crack-tip blunting of the macrocrack are included in the simulations; as dislocations are emitted, the microcrack is kept at a constant distance from the blunted tip of the macrocrack. The brittle-ductile transition (BDT) curve is obtained by simulating the fracture toughness at various temperatures. A consideration of the effects of blunting is found to be critical in predicting the sharp upturn of the BDT curve. The obtained results are compared with existing experimental data and are found to be in reasonable agreement. This article is based on a presentation made in the symposium “Computational Aspects of Mechanical Properties of Materials,” which occurred at the 2005 TMS Annual Meeting, February 13–17, 2005, in San Francisco, CA, under the auspices of the MPMD-Computational Materials Science & Engineering (Jt. ASM-MSCTS) Committee.     

Atomistic Simulations of the Plasticity Behavior of Shock-Induced Polycrystalline Nickel

Shock loading single crystalline nickel creates a defected nanostructure dominated by stacking faults and twins. This transformation is caused by a complex interplay between the incident waves, the waves reflected from sample-free surfaces, and the interference between reflected waves. The plasticity behavior of this shock-induced defected nickel was studied using molecular dynamics (MD) simulations. Compared to a perfect single-crystal nickel sample of the same size, the twinned sample has significantly less yield stress in compression, a slightly lower yield stress in tension, and a yield stress about 30 pct higher in shear. Importantly, our simulations reveal the underlying atomistic mechanisms of dislocation nucleation and twin growth. We observe that while strengthening under shear loading involves lattice dislocations cutting through twins, weakening arises from nucleation of dislocations on the twins under tensile and compressive loading. Also, we have discovered precursors to dislocation loop nucleation in these simulations. This article is based on a presentation made in the symposium entitled “Dynamic Behavior of Materials,” which occurred during the TMS Annual Meeting and Exhibition, February 25–March 1, 2007 in Orlando, Florida, under the auspices of The Minerals, Metals and Materials Society, TMS Structural Materials Division, and TMS/ASM Mechanical Behavior of Materials Committee.     

Intergranular fracture by slip/grain boundary interaction

The role of interaction between slip dislocations and [110] tilt boundaries in crack nucleation has been analyzed for the face-centered cubic (fcc) and L1₂ structures. Dislocation absorption into the boundaries is orientation-dependent according to the elastic anisotropy. Whenthe transfer of slip across the boundary is impeded, cleavage fracture is predicted on the (1ˉ11) plane. Intergranular fracture can be initiated when a symmetric double pileup of primary slip dislocations from both sides of the boundary occurs simultaneously. The available experimental data on slip/grain boundary (GB) interaction and intergranular fracture are in good agreement with the present predictions. This paper is based on a presentation made in the symposium “Interface Science and Engineering“ presented during the 1988 World Materials Congress and the TMS Fall Meeting, Chicago, IL, September 26–29, 1988, under the auspices of the ASM-MSD Surfaces and Interfaces Committee and the TMS Electronic Device Materials Committee.     

A mechanism for transgranular stress-corrosion cracking

A model is proposed to explain transgranular-stress corrosion cracking (T-SCC) in face-centered cubic (fcc) materials. Crack propagation is shown to be anisotropic, in that growth near {110} < 001> is discontinuous due to crack arrest by dislocation blunting whereas growth away from this growth orientation is continuous. For the former case, renucleation of arrested cracks involves active dissolution of shear bands at the crack tip, which changes the stress state at Lomer-Cottrell locks, causing them to fail by cleavage. Once the crack is nucleated, its instantaneous macroscopic crack-growth velocity is considered to be comprised of multiple nucleation of microcracks with intervening arrests. This microcracking results from the interaction of the stress fields from neighboring cracks which are forming simultaneously, the crack-opening constraint due to ligaments which act as “bridges” behind the crack front, and the localized dissolution at the microcrack tip which affectsK _IC and leads to the “cobblestone” appearance. Experimental evidence and theoretical considerations are presented to support the model. The system studied was Cu-25 at. pct Au in 0.6 M NaCl solution at potentials between 300 and 400 mV (sce), which precludes hydrogen embrittlement. This article is based on a presentation made in the symposium “Quasi-Brittle Fracture” presented during the TMS fall meeting, Cincinnati, OH, October 21–24, 1991, under the auspices of the TMS Mechanical Metallurgy Committee and the ASM/MSD Flow and Fracture Committee.     

Determination of dislocation densities in HCP metals from X-ray diffraction line-broadening analysis

X-Ray diffraction (XRD) line-broadening analysis has been performed on highly textured Zr-2.5Nb specimens which had been deformed in tensile tests to produce well-controlled dislocation structures. An iterative deconvolution method has been applied to extract the broadening function for the material, using as standards, a Zr single crystal and a 0 pct deformed specimen. In both cases, for specific tensile tests, a significant contribution to the basal line broadening was observed, which was clearly not directly related to the dislocation structure generated by the deformation, i.e., so-called c-component dislocations having a component of their Burgers vectors perpendicular to the basal plane. Calculations showed that the extent of basal line broadening cannot be attributed to the secondary effect of strain from a-type dislocations, i.e., dislocations with Burgers vectors parallel with the basal plane. It is concluded that most of the line broadening observed was the result of intergranular strain distributions. These distributions are most prominent for grains oriented with their c-axes perpendicular to the tensile-deformation axis and resulted in basal-plane line broadening even when there were few, if any, c-component dislocations present. This article is based on a presentation made in the symposium entitled “Defect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science and Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

Determination of dislocation densities in HCP metals from X-ray diffraction line-broadening analysis

X-Ray diffraction (XRD) line-broadening analysis has been performed on highly textured Zr-2.5Nb specimens which had been deformed in tensile tests to produce well-controlled dislocation structures. An iterative deconvolution method has been applied to extract the broadening function for the material, using as standards, a Zr single crystal and a 0 pct deformed specimen. In both cases, for specific tensile tests, a significant contribution to the basal line braodening was observed, which was clearly not directly related to the dislocation structure generated by the deformation, i.e., so-called c-component dislocations having a component of their Burgers vectors perpendicular to the basal plane. Calculations showed that the extent of basal line broadening cannot be attributed to the secondary effect of strain from a-type dislocations, i.e., dislocations with Burgers vectors parallel with the basal plane. It is concluded that most of the line broadening observed was the result of intergranular strain distributions. These distributions are most prominent for grains oriented with their c-axes perpendicular to the tensile-deformation axis and resulted in basal-plane line broadening even when there were few, if any, c-component dislocations present. This article is based on a presentation made in the symposium entitled “Defect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science and Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

A method of measuring stored energy macroscopically using statistically stored dislocations in commercial purity aluminum

Stored energy from plastic deformation in rolled aluminum has been quantified with both macroscopic and microscopic methods. Differential scanning calorimetry (DSC) and Microhardness tests were used to determine a value for stored energy based on energy released during recrystallization and resistance to plastic flow from the accumulated dislocation content, respectively. For a value of stored energy based only on geometrically necessary dislocations, orientation imaging microscopy (OIM) within a scanning electron microscope (SEM) was used and supported by transmission electron microscopy (TEM) observation of subgrain cell structure. A value for the average misorientation angle that could be associated with the TEM was obtained from the OIM data. The values of stored energy derived from the various analyses were found to be similar with slight overestimation from the OIM technique. Thus, the difference between the macroscopic and microscopic methods represented the statistically stored dislocations. This article is based on a presentation made in the symposium entitled “Processing and Properties of Structural Materials,” which occurred during the Fall TMS meeting in Chicago, Illinois, November 9–12, 2003, under the auspices of the Structural Materials Committee.     

Microstructural evolution during creep of single-phase gamma TiAl

Mechanical experiments and transmission electron microscope (TEM) observations indicate that single-phase γ-TiAl does exhibit primary, secondary, and inverse creep, but not steady-state creep. Constant stress creep tests of γ-Ti-51Al-2Mn conducted at 550 °C, 597 °C, and 703 °C have been interrupted at different stages in the creep process. The TEM observations of these specimens were used to document the microstructural evolution that occurs during creep. Superdislocation motion was activated and subsequently exhausted during primary creep. Ordinary dislocations were observed to be pinned during primary creep, but with time, these dislocations began to bow past their pinning points. The extended region of inverse creep has been related to the bowing and multiplication of these ordinary dislocations. Quantitative measurements of dislocation density were performed, and while the density of superdislocations remained constant, the density of ordinary dislocations increased by an order of magnitude during the life of a creep test. The acceleration in the creep rate has been related to this increase in the density of ordinary dislocations, but the change in dislocation density was not high enough to account for the increase in the creep rate. This suggests that both the mobility and density of ordinary dislocations increase as creep progresses. This article is based on a presentation made in the symposium “Fundamentals of Gamma Titanium Aluminides,” presented at the TMS Annual Meeting, February 10–12, 1997, Orlando, Florida, under the auspices of the ASM/MSD Flow & Fracture and Phase Transformations Committees.     

Atomic-scale modeling of dislocations and related properties in the hexagonal-close-packed metals

Metals with the hcp crystal structure have a wide variety of mechanical and physical properties, and understanding the links between atomic processes, microstructure, and properties can open the way for new applications. Computer modeling can provide much of the information required. This article reviews recent progress in atomic-scale computer simulation in three important areas. The first is the core structure of dislocations responsible for the primary slip modes, where modeling has revealed the variety of core states that can arise in pure, elemental metals and ordered alloys. While most research has successfully employed many-body, central-force interatomic potentials, they are inadequate for metals which have an unfilled d-electron band, such as α-Ti and α-Zr, and the resulting noncentral character of the atomic bonding is shown to have subtle yet significant effects on dislocation properties. Deformation twinning is an important process in plasticity of the hcp metals, and modeling has been used to investigate the factors that control the structure and mobility of twinning dislocations. Furthermore, simulation shows that twinning dislocations are actually generated, in some cases, following the interaction of crystal dislocations with twin boundaries; this can lead to the very mobile boundaries observed experimentally. The final area concerns the nature and properties of the defects created by radiation damage. Computer simulation has been used to determine the number and arrangement of defects produced in primary, displacement-cascade damage in several hcp metals. The number is similar to that found in cubic metals and is considerably smaller than that expected from earlier models. Many self-interstitial atoms cluster in cascades to form highly glissile dislocation loops, and, so, contribute to two-dimensional material transport in damage evolution. This article is based on a presentation made in the symposium entitled “Defect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

Creep crack growth simulation under transient stress fields

The validity of the C -integral for correlation of creep crack growth under transient and steady-state stress fields has been investigated, using FEM analysis. In the steady-state regime, crack growth rates can be correlated withC*, however only for limited amounts of crack extension. When a crack grows in a transient regime no correlation of crack growth is found with any of the conventional crack tip parameters. For both regimes the most relevant parameter is theC-integral values obtained from the near-field region ahead of the crack tip. This paper is based on a presentation made in the symposium “Crack Propagation under Creep and Creep-Fatigue” presented at the TMS/AIME fall meeting in Orlando, FL, in October 1986, under the auspices of the ASM Flow and Fracture Committee.     

Grain-size dependence of the flow stress of Cu from millimeters to nanometers

Data in the literature on the effect of grain size (d) from millimeters to nanometers on the flow stress of Cu are evaluated. Three grain-size regimes are identified: regime I, d>∼10⁻⁶ m; regime II, d ≈10⁻⁸ to 10⁻⁶ m; and regime III, d<∼10⁻⁸ m. Grain-size hardening occurs in regimes I and II; grain-size softening occurs in regime III. The deformation structure in regime I consists of dislocation cells; in regime II, the dislocations are mostly restricted to their slip planes; in regime III, computer simulations indicate that dislocations are absent and that deformation occurs by the shearing of grain-boundary atoms. The transition from regime I to II occurs when the dislocation cell size becomes larger than the grain size, and the transition from regime II to III occurs when the dislocation spacing due to elastic interactions becomes larger than the grain size. The rate-controlling mechanism in regime I is concluded to be the intersection of dislocations; in regime II, it is proposed to be grain-boundary shear promoted by the pileup of dislocations; in regime III, it appears to be grain-boundary shear by the applied stress alone. This article is based on a presentation given in the symposium “Dynamic Deformation: Constitutive Modeling, Grain Size, and Other Effects: In Honor of Prof. Ronald W. Armstrong,” March 2–6, 2003, at the 2003 TMS/ASM Annual Meeting, San Diego, California, under the auspices of the TMS/ASM Joint Mechanical Behavior of Materials Committee.     

Transmission electron microscopy of the crack tip region of fatigued copper single crystals

Dislocation configurations in slices of material cut from the regions ahead of fatigue crack tips in copper single crystals were observed under the transmission electron microscope. Distinct transactions in the configurations of the dislocations occurred at distances approximately equal to 300 μm from the crack tips. The existence of a high density dislocation cell structure in the immediate vicinity of crack tips lends support to the concept of a plastic zone ahead of a fatigue crack tip. The cell diameter decreased the closer the cells were to the tip of a fatigue crack. This paper is based upon a thesis submitted by A. H. PURCELL in partial fulfillment of the requirements of the degree of Master of Science at Northwestern University.     

Brittle to ductile transition in cleavage fracture

In many cleavable solids, cleavage cracks can propagate at steady state by laying down a trail of dislocations emanating from the crack tip and lying on planes inclined to the crack front. These crack-tip initiated dislocations produce shielding at the crack tip that reduces both the crack tip tensile stresses and the shear stresses on the inclined planes. They also blunt the crack. Cleavage cracks, can nevertheless, still propagate under appropriately increased stress intensity conditions to keep the crack tip tensile stress constant. A condition is reached in the propagation of such slightly blunted cracks where a small increment in temperature or a decrement in the crack velocity permits the nucleation of a new set of dislocations that produce additional shielding and blunting which tip the balance against the crack-tip tensile stresses. This results in a transition from brittle cleavage to ductile behavior. The steady state specific plastic work that can just be tolerated by a propagating cleavage crack before it catastrophically blunts is calculated to be only of the order of 10% of the specific surface energy. Although most geometrical details of the dislocation emission process are adequately modeled, the calculated brittle to ductile transition temperatures are found to be more than an order of magnitude higher than those that have been experimentally measured. This discrepancy is a result of the present inadequate methods of modeling activation configurations by considering the dislocation loop radius as the only activation parameter, while proper modeling of such configurations must consider also the Burgers shear displacement of the loop as an activation parameter. Such two parameter analyses, however, require accurate information on interlayer atomic shear resistance profiles for specific crystals which are presently not available. The analysis furnishes ready explanations of the toughening effects of so-called “ductilization” treatments and embrittling effect of aging and dislocation locking, as well as the relatively large difference between the lowest levels of toughness between fracture in polycrystals and in single crystals.