The effects of deformation and pre-heat-treatment on abnormal grain growth in RENÉ 88 superalloy

When an extruded strain-free RENé 88 Ni-base superalloy about 1 to 2 μm in grain diameter is heattreated at 1150°C, abnormal grain growth (AGG) begins after 50 hours. When heat-treated at 1200 °C, AGG occurs at 15 minutes. Some of the grain boundaries are faceted with hill-and-valley structures when observed in transmission electron microscopy (TEM), and the occurrence of AGG is consistent with the boundary step and dislocation mechanism for the migration of singular boundaries with faceted shapes, as observed and proposed in other pure metals and alloys. The dissolution of abundant γ′ precipitates (with a solvus temperature of 1107 °C), which are coherent with the matrix and hence strongly pin the grain boundaries, does not cause AGG during early stages of heat-treatment at 1150 °C. Small deformations drastically alter the AGG behavior. When deformed to 4 pct, AGG begins after heat-treating for 10 minutes at 1150 °C, compared to the apparent incubation time of 50 hours for an undeformed specimen, and very large abnormal grains are produced. With increasing deformation to 6 and 9.2 pct, the abnormal grain size decreases. These results are qualitatively similar to those observed in Cu. This deformation effect on AGG is attributed to the absorption of lattice dislocations in the grain boundaries, which produces nonequilibrium structures that, in turn, can apparently cause rapid boundary migration. When heat-treated at 1200 °C, the largest abnormal grains are found in the specimens deformed to 2 pct. When the initial grain size is increased to about 14 μm by heattreating the extruded alloy at 1150 °C for 30 minutes, similar low deformation effects on AGG are observed. When these specimens are deformed to 10, 13, and 15.2 pct, primary recrystallization occurs during the heat-treatment at 1150 °C, and large abnormal grains are again produced because of the small recrystallized grain size. Therefore, there are two peaks in the grain size vs deformation curve after heat-treating at 1150 °C for 1 hour. A pre-heat-treatment of this alloy at 1050 °C below the solvus temperature of the γ′ phase greatly reduces the size of the abnormal grains obtained in the specimen deformed to 4 pct after the heat-treatment at 1150 °C, probably because some recovery of the dislocations takes place at grain boundaries during the pre-heat-treatment. The deformation effect on AGG observed in this alloy is qualitatively similar to that previously observed in Cu and appears to be consistent with the boundary step and dislocation mechanism for AGG. An erratum to this article is available at .     

Abnormal grain growth in bulk Cu—The dependence on initial grain size and annealing temperature

The dependence of abnormal grain growth (AGG), also termed secondary recrystallization, on annealing temperature in the range between 600 °C and 1050 °C has been observed in pure bulk Cu specimens compressed to various levels between 5 and 75 pct. There is no grain texture after annealing. The average grain size after primary recrystallization, which represents the initial grain size for secondary recrystallization during further annealing, decreases with increasing deformation and is nearly independent of the annealing temperature, in agreement with previous observations. The incubation time for AGG decreases and the number density of abnormally large grains increases with increasing deformation (hence, a decreasing initial grain size) and increasing annealing temperature. At low temperatures, most of the grain boundaries are faceted, with some facet planes probably of singular structures corresponding to cusps in the polar plots of the grain-boundary energy vs the grain-boundary normal. With increasing temperature, the grain boundaries become defaceted and, hence, atomically rough. The observed grain-growth behavior appears to be qualitatively consistent with the movement of faceted grain boundaries by two-dimensional nucleation of boundary steps. The temperature dependence appears to be consistent with roughening of grain boundaries. Before the onset of AGG, stagnant growth of the grains occurs at low rates, probably limited by slow two-dimensional nucleation of boundary steps, and, at low deformations and low annealing temperatures, the stagnant growth persists for 100 hours. The specimens with relatively small initial grain sizes (because of high deformation) show double AGG when annealed at high temperatures.     

The dependence of normal and abnormal grain growth in silver on annealing temperature and atmosphere

When polycrystalline pure Ag specimens are compressed to 40 pct and annealed, there is no noticeable texture in the recrystallized structure, and normal or abnormal grain growth occurs during annealing. When annealed further in low vacuum (10⁻³ to 10⁻⁴ Torr) after the completion of recrystallization, normal grain growth occurs at 920 °C and 800 °C, but abnormal grain growth (AGG) occurs at 700 °C, 600 °C, and 500 °C. When annealed in O₂ atmosphere, normal grain growth occurs at 920 °C and AGG at 800 °C, 700 °C, 600 °C, and 500 °C. At temperatures close to the melting point (960.5 °C), the grain boundaries are expected to be rough at atomic scales and hence have nearly isotropic boundary energy. The normal growth of the grains with such atomically rough boundary structures is consistent with some theoretical analysis and simulation. At low temperatures, the grain boundaries can be faceted with probably singular structures. Because these grain boundaries apparently migrate by the movement of boundary steps, AGG occurs. The observations with optical microscopy indeed indicate that some grain boundaries are faceted at low temperatures and all of them are smoothly curved indicating an atomically rough structure at high temperatures close to the melting point. Although the results are not conclusive, they support the hypothesis that AGG occurs because the faceted singular grain boundaries migrate by the step mechanism.     

Abnormal growth of faceted (WC) grains in a (Co) liquid matrix

If the grains dispersed in a liquid matrix are spherical, their surface atomic structure is expected to be rough (diffuse), and their coarsening has been observed to be controlled by diffusion in the matrix. They do not, furthermore, undergo abnormal growth. On the other hand, in some compound material systems, the grains in liquid matrices are faceted and often show abnormal coarsening behavior. Their faceted surface planes are expected to be singular (atomically flat) and therefore grow by a defect-assisted process and two-dimensional (2-D) nucleation. Contrary to the usual coarsening the-ories, their growth velocity is not linearly dependent on the driving force arising from the grain size difference. If the growth of the faceted grains occurs by 2-D nucleation, the rate is expected to increase abruptly at a critical supersaturation, as has been observed in crystal growth in melts and solutions. It is proposed that this growth mechanism leads to the abnormal grain coarsening. The 2-D nucleation theory predicts that there is a threshold initial grain size for the abnormal grain growth (AGG), and the propensity for AGG will increase with the heat-treatment temperature. The AGG behavior will also vary with the defects in the grains. These predictions are qualitatively confirmed in the sintered WC-Co alloy prepared from fine (0.85-Μm) and coarse (5.48-Μm) WC powders and their mixtures. The observed dependence of the AGG behavior on the sintering temperature and the milling of the WC powder is also qualitatively consistent with the predicted behavior.     

Static recrystallization mechanisms in a coarse-grained Nb-microalloyed austenite

Static recrystallization mechanisms have been studied in a coarse-grained Nb microalloyed austenite. An austenite with a coarse grain size of 800 μm, typical of thin slab casting processes, has been deformed in torsion at a temperature of 1100 °C. After deformation, the specimens have been held for different times at this high temperature and then water quenched. The microstructural changes occurring during static recrystallization were characterized by metallographic evaluation. It has been observed that new recrystallized grains nucleate preferentially on parent austenite grain boundaries and tend to form in clusters. Once all the boundaries have been consumed, intragranular nucleation is actived at late stages of recrystallization. Clustered nucleation allows impingement to take place early during the recrystallization process, favoring grain-coarsening phenomena to occur behind the recrystallization front, which is denoted by the significant reduction in the number of grains per unit volume observed during early stages of recrystallization. Static recrystallization proceeds heterogeneously, as a result of a nonuniform distribution of stored energy in the deformed material. A continuous decrease of the average migration rate of the recrystallization front is observed, which can be ascribed to the reduction of the driving force for migration as recrystallization advances.     

Ferrite recrystallization and austenite formation in cold-rolled intercritically annealed steel

The recrystallization of ferrite and austenite formation during intercritical annealing were studied in a 0.08C-1.45Mn-0.21Si steel by light and transmission electron microscopy. Normalized specimens were cold rolled 25 and 50 pct and annealed between 650 °C and 760 °C. Recrystallization of the 50 pct deformed ferrite was complete within 30 seconds at 760 °C. Austenite formation initiated concurrently with the ferrite recrystallization and continued beyond complete recrystallization of the ferrite matrix. The recrystallization of the deformed ferrite and the spheroidization of the cementite in the deformed pearlite strongly influence the formation and distribution of austenite produced by intercritical annealing. Austenite forms first at the grain boundaries of unrecrystallized and elongated ferrite grains and the spheroidized cementite colonies associated with ferrite grain boundaries. Spheroidized cementite particles dispersed within recrystallized ferrite grains by deformation and annealing phenomena were the sites for later austenite formation.     

Orientation selective recrystallization of nonoriented electrical steels

A nonoriented electrical steel that was commercially hot rolled and then given a 70 pct cold reduction on a laboratory mill was annealed at 680 °C for 6 minutes. The sheet was then submitted to a second rolling reduction of 5.2 pct, followed in turn by a second annealing at 730 °C for various times. The textures were measured after the first and second recrystallization treatments and analyzed using a nucleation and growth model. In the model, the nucleus orientation distribution function is first calculated by assessing the nucleation probability for each deformed matrix orientation. The nucleation texture is then transformed into the recrystallization texture by means of an appropriate growth criterion. The calculations indicate that the annealing texture of the conventionally rolled (70 pct reduction) sheet can be accounted for on the basis of random nucleation followed by selective growth. The latter is characterized by the following physical features: (a) the low mobility of low angle grain boundaries, (b) the enhanced mobility of {110} plane matching boundaries, and (c) variant selection of the (110) plane that carries the largest amount of slip during deformation. The computer simulations also show that low stored energy nucleation is favored in the lightly rolled sheet. These nuclei grow into the matrix by a selection mechanism that involves the increased mobility of 219a and 233a (110) coincident site lattice (CSL) boundaries.     

Early Stages of Recrystallization in Equal-Channel Angular Pressing (ECAP)-Deformed AA3104 Alloy Investigated Using Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM) Orientation Mappings

Early stages of recrystallization in alloys containing complex structure of second phase particles are of considerable practical interest. They were observed for the AA3104 alloy in which large particles occur against the background of randomly distributed dispersoids. The samples were deformed by equal channel angular pressing and then slightly annealed to obtain the state of partial recrystallization. The highly deformed alloy contained a structure of flat grains with the spacing between high-angle grain boundaries ranging from 100 nm to 1 ??m. On annealing, the structure coarsened and got transformed into nearly equiaxed grains by both discontinuous and continuous recrystallization. The nucleation of new grains was observed in statically recrystallized bulk samples using scanning electron microscopy, and during in-situ recrystallization in a transmission electron microscope. Special attention was paid to the nucleation of new grains in areas close to large second phase particles, where a relatively high stored energy was expected to stimulate nucleation. A particular role in the rise of nuclei is attributed to migration of low angle boundaries. During recrystallization at 623 K (350?°C), in most of the observed cases, the growth of grains occurred by coalescence of neighbouring cells and by migration of high-angle grain boundaries. These processes led to nearly equiaxed grains of similar size. Orientation mappings showed that although orientations of new grains were widely scattered, they were not completely random.     

A Comparative Study on the Static Recrystallization Behavior of Cold-Rolled Mg-3Al-1Zn Alloy Stimulated by Electropulse Treatment and Conventional Heat Treatment

Cold-rolled AZ31 Mg alloy strips, with a reduction of 33 pct, were subjected to electropulse treatment (EPT) and conventional heat treatment (HT) to evaluate the respective influences of electropulses and temperature on the recrystallization behavior of AZ31. The highest measured temperature during the EPT (543 K) was used in HT. The electron backscattered diffraction results demonstrated that the EPT-stimulated recrystallization was completed within 8 seconds, whereas for HT, recrystallization was still far from completion even after 240 seconds. It was found that both the nucleation and grain growth of these two processes were totally different. In the EPT samples, nucleation tended to occur preferentially near extension twin boundaries and grain boundaries by continuous recrystallization, whereas in the HT samples, nucleation occurred mainly by grain boundaries bulging via discontinuous recrystallization. As grain growth proceeded, the texture intensities of the EPT samples decreased gradually and finally evolved into an obvious transverse-direction-split texture. This is likely attributable to the impact of electropulses on the boundary energy and the contribution of nonbasal dislocations; however, the basal-type textures of the HT samples were notably strengthened, which is associated with a 30 deg〈0001〉 orientation with respect to the deformed texture.     

Recrystallization and formation of austenite in deformed lath martensitic structure of low carbon steels

The effect of prior deformation on the processes of tempering and austenitizing of lath martensite was studied by using low carbon steels. The recrystallization of as-quenched lath martensite was not observed on tempering while the deformed lath martensite easily recrystallized. The behavior of austenite formation in deformed specimens was different from that in as-quenched specimens because of the recrystallization of deformed lath martensite. The austenitizing behavior (and thus the austenite grain size) in deformed specimens was controlled by the competition of austenite formation with the recrystallization of lath martensite. In the case of as-quenched (non-deformed) lath martensite, the austenite particles were formed preferentially at prior austenite grain boundaries and then formed within the austenite grains mainly along the packet, block, and lath boundaries. On the other hand, in the case of lightly deformed (30 to 50 pct) lath martensite, the recrystallization of the matrix rapidly progressed prior to the formation of austenite, and the austenite particles were formed mainly at the boundaries of fairly fine recrystallized ferrite grains. When the lath martensite was heavily deformed (75 to 84 pct), the austenite formation proceeded almost simultaneously with the recrystallization of lath martensite. In such a situation, very fine austenite grain structure was obtained most effectively.     

Computer simulation of recrystallization—III. Influence of a dispersion of fine particles

Two-dimensional Monte Carlo simulations of recrystallization have been carried out in the presence of incoherent and immobile particles for a range of different particle fractions, a range of stored energies and a range of densities of potential nuclei (embryos). For stored energies greater than a critical value (H/J > 1) the recrystallization front can readily pass the particles leading to a random density of particles on the front and a negligible influence of particles on the recrystallization kinetics. At lower stored energies the particles pin the recrystallization front leading to incomplete recrystallization. However at very low particle fractions, when the new grain has grown much larger than the matrix grains, before meeting any particles, the new grains can complete the consumption of the deformed grains giving complete “recrystallization” by a process that appears to be similar to abnormal grain growth. Particles are, as reported previously, very effective at pinning grain boundaries, both of the deformed and recrystallized grains, when boundaries migrate under essentially the driving force of boundary energy alone. Such boundaries show a density of particles that rises rapidly from the random value found at the start of the simulation. As a consequence, particles very strongly inhibit normal grain growth after recrystallization. Such growth can only occur if the as-recrystallized grain size is less than the limiting grain size seen in the absence of recrystallization. Under these circumstances a small increment of grain growth occurs until the grain boundaries once again acquire a higher than random density of particles.     

Initial stages of recrystallization in aluminum of commercial purity

In commercial aluminum with a purity of 99.4 pct, the formation and growth of recrystallization nuclei were studied by techniques such asin-situ annealing in a high voltage electron microscope, transmission electron microscopy and light microscopy. Sample parameters were the initial grain size (370 and 19 microns) and the degree of deformation (50 and 90 pct reduction in thickness by cold-rolling). It was found that the initial grain boundaries and high angle boundaries within the original grains are preferential sites for recrystallization nuclei, and that the effect of such sites is enhanced by the FeAl₃ particles present in the commercial aluminum as impurities. The nucleation temperatures determined by high voltage electron microscopy and transmission electron microscopy decrease markedly when the initial grain size is decreased both after 50 and 90 pct cold rolling; a less pronounced temperature decrease is obtained by increasing the degree of deformation. The size of the recrystallization nuclei, the recrystallization temperature and the recrystallized grain size are reported for the four sample states, and finally the structural and kinetic observations are discussed.     

Effect of gas bubbles on recrystallization of tungsten

Tungsten sheet deposited from WF₆ vapor was rolled at temperatures up to 1000°C and reductions in thickness up to 95 pct. Recovery and recrystallization were studied using hardness measurements and transmission electron microscopy. The temperatures for the onset of recrystallization during 1 hr anneals were 1600°C or more. Recrystallization was sluggish, requiring temperatures of 2100°C or more for completion. Transmission electron microscopy revealed that many small (< 500Å diam) gas bubbles were present on dislocations and recovered subgrain boundaries and caused scalloping of these boundaries and the advancing recrystallized grain boundaries. Coarsening of the bubbles to a relatively few larger bubbles permitted complete recrystallization. No solid inclusions were found. Annealing treatments that coarsened the gas bubbles prior to rolling caused accelerated recovery and complete recrystallization was observed at approximately 1200°C. It is concluded that small gas bubbles were the cause of the original high recrystallization temperatures and also influenced the size and shape of the recrystallized grains.     

On deformation-induced continuous recrystallization in a superplastic AlLiCuMgZr alloy

The microstructural changes of a warm rolled AlLi alloy occurring during static annealing and superplastic deformation at 515°C were studied by means of transmission electron microscopy. Deformation induces a continuous recrystallization with a rapid subgrain growth and a rapid increase in boundary misorientations. The higher strain rate results in a faster subgrain growth and a finer recrystallized grain size. The increasing rate of boundary misorientations and the strain at which the average misorientation reaches about 20° increase with increasing strain rate. The increase in boundary misorientations is proportional to the subgrain growth during the whole static annealing process. Deformation results in a more rapid increase in boundary misorientation with subgrain size than static annealing. Dislocation gliding plays an important role before the formation of high angle grain boundaries during superplastic deformation. The absorption of dislocations into subgrain boundaries results in a more rapid increase in boundary misorientation during deformation. Thus, the mechanism of the deformation-induced continuous recrystallization is suggested to be the generation of dislocations in grains and the absorption of gliding dislocations into subgrain boundaries.     

Nb(C,N) Precipitation and Austenite Recrystallization

The influence of boron addition, amount of deformation, and solution heat-treatment temperature on the precipitation and recrystallization behaviors of a family of high-strength low-alloy (HSLA) steels was studied. A stress relaxation technique was employed to detect the occurrence of austenite recrystallization and to determine the precipitation start(P _s) and finish(P _f) times. After preheating to 1100 °C or 1200 °C for 30 minutes, the specimens were cooled to test temperatures between 800 °C and 1000 °C. They were subsequently deformed to true strains of 5 or 25 pct and subjected to stress relaxation. The advent of recrystallization produced a sharp increase in relaxation rate, while the occurrence of carbonitride precipitation led to the appear- ance of a stress plateau. The results indicate that the presence of boron (1) accelerates carbo- nitride precipitation and (2) retards austenite recrystallization when present in combination with Nb. The precipitation-time-temperature (PTT) diagrams determined in this investigation are C-shaped for both the B-free as well as the B-modified Nb steels. These data were analyzed in terms of the classical theory of nucleation, on the basis of which it is demonstrated that the acceleration of the nucleation kinetics of precipitation can be attributed to the segregation of boron and of boron-vacancy complexes to dislocations and grain boundaries, as well as to the faster diffusion of Nb in the presence of boron.     

Formation of Cube and Goss Texture After Primary Recrystallization in Electrical Steels

Recrystallization texture evolution in 2.1 pct Si electrical steel sheets was investigated with X-ray diffraction and electron backscattered diffraction (EBSD) analysis. The beneficial Cube ({100} 〈001〉 ) and Goss ({110} 〈001〉 ) components dominate primary recrystallization texture after final annealing at 1123 K (850 °C). It is found that Cube grains nucleate within {114} 〈481〉  deformed grains, and the nucleation of Cube grains is later than that of Goss grains during recrystallization. The formation of strong Cube and Goss texture is explained in terms of both the oriented nucleation and selective growth mechanisms.     

Recovery,recrystallization and grain growth in 0.03 C, 1.8 Si, 0.3 Al steels with and without addition of 0.05% of antimony

The investigation was carried out on two laboratory steels elaborated from identical base materials, one with 0.05% Sb and an industrial steel as-delivered and after decarburisation. Cold rolled 0.5 mm sheets were prepared by laboratory rolling and investigated after annealing in temperature range 550 to 800°C for 0.5 to 60 m. Antimony has no effect on recovery in temperature range 550 to 625°C. Rare recrystallization nuclei were found at grain boundaries only in the decarburized steel, in all other steels nucleii appeared and grew only inside of deformed grains. At recrystallization finished grains were coarser in antimony and decarburized steels. The explanation is that less numerous nuclei grew for a longer time in the deformed matrix. The size of recrystallized grains was proportional to the square root of the annealing time. Recrystallized grain growth was faster in antimony-free laboratory and in decarburized industrial steels, while the growth activation energy is very similar in antimony and decarburized steels. The growth activation energy was greater in antimony-free laboratory and in as-delivered industrial steels. It was similar to the activation energy of part of the recovery and near to the activation energy of iron self-diffusion in ferrite. No difference was found in grain growth topology between the antimony and the comparative laboratory steels. Results indicate a similar effect of the presence of 0.05% of antimony and the decreased carbon content in steel.     

Grain boundary faceting and abnormal grain growth in nickel

A correlation between grain boundary faceting and abnormal grain growth has been observed in recrystallized polycrystalline Ni at varying annealing temperatures, with or without C added. Carburized Ni specimens deformed to 50 pct show faceted grain boundaries and abnormal grain growth when annealed at temperatures below 0.7 T _m, where T _m is the melting point of Ni in absolute scale. When annealed at or above 0.7 T _m, the grain boundaries are smoothly curved and, therefore, have a rough structure, and normal grain growth is observed. In the specimens annealed in vacuum without carburization, all grain boundaries are faceted at 0.55 T _m, and some of them become defaceted at higher temperatures. The specimens annealed in vacuum at temperatures between 0.55 and 0.95 T _m show abnormal grain growth. When the grain boundaries have a rough structure and are, therefore, nearly isotropic, normal grain growth is indeed expected, as shown by the simulation and analytical treatment. When all or a fraction of the grain boundaries are faceted, with the facet planes corresponding to the singular cusp directions in the variation of the boundary energy against the inclination angle, abnormal grain growth can occur either because some grain boundary junctions become immobile due to a torque effect, or the growth occurs by a step mechanism.     

Quantifying recrystallization nucleation and growth kinetics of cold-worked copper by microstructural analysis

Microstructural evolution data describing the recrystallization of cold-worked copper at 394 K (121 °C) were obtained by quantitative metallography using scanning electron microscopy and electron backscattered pattern analysis. Using the microstructural path method (MPM), a new analytical representation of the microstructure was devised that emulated all the measurements and successfully explained why simpler representations failed to adequately describe the kinetics of recrystallization in copper. Saturation of preferentially located nucleation sites such as at deformation bands, grain boundaries, etc., where recrystallized grains may cluster in planar arrays before the deformed volume is completely consumed, and time-dependent growth rates matched fully the kinetic behavior of copper during recrystallization. The kinetic behavior of individual texture components (random and cube + cube twin) was also delineated, experimentally and analytically. Precise matching of the analytical representation of the microstructure to experiment allowed calculation of nucleation and growth parameters. These showed that the cube + cube twin grains nucleated at a faster rate than the random grains, that site saturation occurred sooner for the cube + cube twin grains, and that cube + cube twin grains grew at rates about 1.5 times faster than the random grains. The calculations suggested that as recrystallization approached completion, the number of random grains slightly outnumbered the cube + cube twin grains.     

Effect of residual magnesium content on thermal fatigue cracking behavior of high-silicon spheroidal graphite cast iron

This study investigates the thermal fatigue cracking behavior of high-silicon spheroidal graphite (SG) cast iron. Irons with different residual magnesium contents ranging from 0.038 to 0.066 wt pct are obtained by controlling the amount of spheroidizer. The repeated heating/cooling test is performed under cyclic heating in various temperatures ranging from 650 °C to 800 °C. Experimental results indicate that the thermal fatigue cracking resistance of high-silicon SG cast iron decreases with increasing residual magnesium content. The shortest period for crack initiation and the largest crack propagation rate of the specimens containing 0.054 and 0.060 wt pct residual magnesium contents are associated with heating temperatures of 700 °C and 750 °C. Heating temperatures outside this range can enhance the resistance to thermal fatigue crack initiation and propagation. When thermal fatigue cracking occurs, the cracks always initiate at the surface of the specimen. The major path of crack propagation is generally along the eutectic cell-wall region among the ferrite grain boundaries, which is the location of MgO inclusions agglomerating together. On the other hand, dynamic recrystallization of ferrite grains occurs when the thermal cycle exceeds a certain number after testing at 800 °C. Besides, dynamic recrystallization of the ferrite matrix suppresses the initiation and propagation of thermal fatigue cracking.