Volume 09 - Metallography and Microstructures Part 16 ppt

Fig. 8 Annealed at 400 °C (750 °F) for 1280 min. Reduced dislocation density and random arrays of dislocations are evident. Fig. 10 Annealed at 800 °C (1470 °F) for 5 min. Increased average diameter of the subgrains is due to subgrain growth. Fig. 7 High density of dislocations and no well- defined cell structure is revealed in the as-rolled condition. Fig. 9 Annealed at 600 °C (1110 °F) for 1280 min. Well- defined subgrains resulting from polygonization are shown. Effect of annealing time and temperature on the microstructure of an Fe-3Si single crystal, cold rolled 80% in the (001)[110] orientation. Thin-foil TEM specimens prepared parallel to the rolling plane. All at 17,2000× When the microstructure of a heavily rolled crystal is revealed using thin-foil specimens parallel to the cross sections of the strip, the thin, ribbon-like deformation cells are readily observed. Figure 11 shows the dislocation substructure of an as-deformed iron crystal cold rolled in the (111)[110] orientation to 70% reduction (see also Fig. 4 for a much finer microstructure in a heavily rolled polycrystalline stainless steel). During recovery, the thickness of the ribbon-like subgrains increases, as shown in Fig. 12. Subgrain growth at these early stages cannot be clearly observed when thin-foil specimens parallel to the rolling plane are used for transmission electron microscopy examination. As mentioned earlier, this is because a clearly defined subgrain structure can be observed in a thin-foil specimen parallel to the rolling plane only when the thickness of the subgrains exceeds that of the foil. Fig. 11 Fig. 12 Electrolytic iron single crystal, cold rolled 70% in the (111)[ 1 10] orientation. Fig. 11: thin, ribbon-like cells stacked up in the thickness dimension of the as-rolled crystal. Fig. 12: annealed at 550 °C (1020 °F) for 20 min. Increased cell thickness resulting from subgrain growth. Thin-foil TEM specimens prepared parallel to the transverse cross section. Both at 11,000×. (Ref 4) References cited in this section 2. R.R. Eggleston, J. Appl. Phys., Vol 23, 1952, p 1400 3. J.T. Michalak and H.W. Paxton, Trans. AIME, Vol 221, 1961, p 850 4. B.B. Rath and H. Hu, in Proceedings of the 31st Annual Meeting of the Electron Microscopy Society of America, San Francisco Press, 1973, p 160 Recrystallization Following recovery, recrystallization (or primary recrystallization) occurs by the nucleation and growth of new grains, which are essentially strain-free, at the expense of the polygonized matrix. During incubation, stable nuclei are formed by the coalescence of subgrains that leads to the formation of high-angle boundaries. From that time on, subsequent growth of new grains can proceed rapidly, because of the high mobility of the high-angle boundaries. The rate of recrystallization later decreases toward completion as concurrent recovery of the matrix occurs and more of the new grains impinge upon each other. Accordingly, isothermal recrystallization curves are typically sigmoidal (see Fig. 24 and 30). Because recrystallization is accomplished by high-angle boundary migration, a large change in the texture occurs. Sufficient deformation and a sufficiently high temperature of annealing are required to initiate recrystallization following recovery. With a low degree of deformation and a low annealing temperature, the specimen may recover only without the occurrence of recrystallization. In situ recrystallization, or complete softening without the nucleation and growth of new grains at the expense of the polygonized matrix, is a process of recovery, not recrystallization, because it does not involve high-angle boundary migration. Consequently, there is no essential change in texture following in situ recrystallization. Nucleation Sites. Because of the highly nonhomogeneous microstructure of a plastically cold-worked metal, recrystallization nuclei are formed at preferred sites. Examples of preferred nucleation sites include the original grain boundaries; the boundaries between deformation bands within a crystal or grain; the intersections of mechanical twins, such as Neumann bands in body-centered cubic crystals; the distorted twin-band boundaries; and the regions of shear bands. Limited recrystallization may also occur by the growth of grains nucleated at large and hard inclusion particles. In general, preferred nucleation sites are regions of relatively small volume where the lattice is highly distorted (having high lattice curvature). In such regions, the dimension of the substructure is fine, and the orientation gradient is high. Therefore, the critical size for a stable nucleus to form in these regions is relatively small and so can be attained more readily. Furthermore, the nucleus needs only to grow through a relatively short distance to form a high-angle boundary with the matrix. Figure 13 shows recrystallized grains formed in the boundary region between two main deformation bands in a crystal of Fe3Si that was cold rolled 80% in the (001)[100] orientation, then annealed at 600 °C (1110 °F) for 25 min. What appears as a thin-line boundary between main deformation bands in an optical micrograph (Fig. 13) actually contains a group of narrow, elongated, microband segments, among which nucleation occurs by the coalescence of the segments into a recrystallized grain (Fig. 14). Nucleation in such microband regions is also termed transition band nucleation, because the large orientation difference between the main deformation bands is accommodated in small steps by the microband segments. In heavily deformed polycrystalline specimens, such transition regions must exist between different deformation texture components, but they may not be as clearly defined and readily identified as in the similarly deformed specific single-crystal specimens. Fig. 13 Fe- 3Si single crystal, cold rolled 80% in the (001)[100] orientation and annealed at 600 °C (1110 °F) for 25 min. Optical micrograph shows recrystallized grains formed at boundaries (microband region or transition bands) between the main deformation bands. See also Fig. 16. 5% nital. 400× Fig. 14 Fe- 3Si single crystal, cold rolled 80% in the(001)[100] orientation and annealed at 600 °C (1110 °F) for 125 min. Transmission electron micrograph showing a recrystall ized grain grown from the microband region (transition bands). Thin-foil specimen prepared parallel to the rolling plane. Compare with Fig. 15. 14,740× Figures 15 and 16 show the nucleation of recrystallized grains in heavily rolled polycrystalline copper by the coalescence of subgrains in the microband regions. These micrographs, which were obtained from thin-foil specimens prepared parallel to the cross section of the sheet, show the evolution of the microstructure in nucleation. When thin-foil specimens prepared parallel to the rolling plane of the heavily rolled sheet are used for nucleation studies, the characteristics of the nucleation site cannot be defined with certainty (Fig. 17 and 18). Fig. 15 Fig. 16 Electrolytic copper, cold rolled 99.5%. Fig. 15: annealed at 100 °C (212 °F) for 625 min. Recrystallization nuclei formed among microbands are shown. 17,100×. Fig. 16: annealed at 100 °C (212 °F) for 25 min. Recrystallization nuclei formed among microbands by subgrain coalescence are shown. 34,200×. Both thin-foil TEM specimens prepared parallel to the transverse section Fig. 17 Low-carbon steel, cold rolled 70% and annealed at 450 °C (840 °F) for 260 h and 42 min. Well- developed recrystallized grains and recrystallization nuclei during their formation by subgrain coalescence in the recovered matrix still exhibit a "messy" substructure. Thin- foil TEM specimen prepared parallel to the rolling plane. 7,020× Fig. 18 Type 304L stainless steel, cold rolled 90% at 25 °C (75 °F) and annealed at 600 °C (1110 °F) for 1 h. Early recrystallized grains with annealing twins in a highly "messy" matrix. Thin- foil TEM specimen prepared parallel to the rolling plane. 21,600×. (Ref 5) In moderately deformed samples with relatively coarse initial grains, the microstructure near the grain boundaries and the evolution of the microstructure during nucleation can be studied in considerable detail, even when thin-foil specimens parallel to the rolling plane are used for transmission electron microscopy examinations. Figure 19 shows the grain- boundary bands observed adjacent to an initial grain boundary in commercial-purity aluminum that was cold rolled 50%. The cumulative misorientations across the bands (16.5°), as shown in the inset, indicate similarity in feature between these grain-boundary bands and the transition bands described earlier. These grain-boundary bands obviously would not form at every grain boundary, but would depend on the relative orientations of the two adjacent grains. Fig. 19 Fine-grained commercial-purity aluminum, cold rolled 50%. A 9-μ m wide grain boundary band consisting of elongated subgrains that was developed along an initial grain boundary mark ed by arrows. The inset shows the misorientations regarding the grain interior as a function of the distance from the grain boundary. Thin-foil TEM specimen prepared parallel to the rolling plane. 7,300×. (Ref 6) Grain-boundary nucleation by the "bulging out" of a section of an initial boundary from the region of a low dislocation content into a region of high dislocation content is frequently observed in large-grained materials deformed at low and medium strains. This bulging mechanism of nucleation for recrystallization is a consequence of the strain-induced boundary migration. Figure 20 shows a recrystallization nucleus that has formed by straddling a grain boundary in a coarse-grained aluminum that was cold rolled 30% and annealed at 320 °C (610 °F) for 30 min. Such grain-boundary nucleation was observed to have three types of structural detail. As shown in Fig. 21, the nucleus may be formed by subgrain growth to the right of the original grain boundary (Fig. 21a), by grain-boundary migration to the right and subgrain growth to the left forming a new high-angle boundary (Fig. 21b), and by grain-boundary migration to the right and subgrain growth to the left but without forming a new high-angle boundary (Fig. 21c). Fig. 20 Coarse-grained commercial- purity aluminum cold rolled 30% and annealed at 320 °C (610 °F) for 30 min. A recrystallization nucleus (denoted A) developed near arrow-marked FeAl 3 particles, and is shown straddling an initial grain boundary (marked by dotted line). Thin- foil TEM specimen prepared parallel to the rolling plane. 3,650×. (Ref 6) Fig. 21 Schematic showing three types of grain- boundary nucleation and the growth of the nucleus (N) at the expense of the polygonized subgrains. See text for detailed explanation. (Ref 6) When a polycrystalline specimen is deformed to a very small strain less than 2 or 3%, for example then annealed at a sufficiently high temperature, recrystallization occurs by strain-induced boundary migration of only a few grains. These few grains grow very large at the expense of the small matrix grains. The maximum level of strain below which such coarsening occurs is commonly termed critical strain. This behavior has been used to grow single crystals in the solid state by the so-called "strain-anneal" technique. Figure 22 shows recrystallized grains nucleated and grown at a large and hard FeAl 3 inclusion particle in 90% cold-rolled aluminum after annealing in the high-voltage electron microscope at 264 °C (507 °F) for 480 s. Unless the volume fraction of the inclusion particles is substantially large, the contribution of particle-nucleated grains constitutes only a small fraction of the total recrystallization volume. From the above discussions on nucleation sites, it is easy to understand that the size of the recrystallized grains, as recrystallization is complete, decreases with increasing deformation, because the number of nuclei increases with increasing deformation. Fig. 22 Fine-grained commercial-purity aluminum, cold rolled 90% and heated in a high- voltage electron microscope at 264 °C (507 °F) for 480 s. Recrystallized grains (denoted by letters) nucleated at a large FeAl 3 particle and grown into the polygonized matrix. Thin- foil TEM specimen prepared parallel to the rolling plane. 2,810×. (Ref 6) Growth of Nucleated Grains. The growth of the newly formed strain-free grains at the expense of the polygonized matrix is accomplished by the migration of high-angle boundaries. Migration proceeds away from the center of boundary curvature. The driving force for recrystallization is the remaining strain energy in the matrix following recovery. This strain energy exists as dislocations mainly in the subgrain boundaries. Therefore, the various factors that influence the mobility of the high-angle boundary or the driving force for its migration will influence the kinetics of recrystallization. For example, impurities, solutes, or fine second-phase particles will inhibit boundary migration; therefore, their presence will retard recrystallization. Figure 23 shows the pinning of a mobile low-angle boundary by a fine alumina (Al 2 O 3 ) particle in an aluminum-alumina specimen during recovery. In connection with the driving force for recrystallization, a fine-subgrained matrix has a higher strain-energy content than does a coarse-subgrained matrix. Accordingly, recrystallization occurs faster in a fine-subgrained matrix than in a coarse-subgrained matrix. During recrystallization, continued recovery may occur in the matrix by subgrain growth, resulting in a reduction of the driving energy for recrystallization and therefore a decrease in the recrystallization rate. From driving energy considerations, it is understandable that the tendency for recrystallization is stronger in heavily deformed than in moderately or lightly deformed specimens. For a given deformation, the finer the original grain size the stronger the tendency for recrystallization. Figure 24 shows such effects in low-carbon steel. Fig. 23 Aluminum-aluminum oxide specimen, cold rolled and annealed. Shown is the pinning of a mobile low- angle boundary by a small Al 2 O 3 particle during a recovery anneal. Thin-foil TEM specimen. 47,000×. (Ref 7) Fig. 24 Effect of penultimate grain size on the recrystallization kinetics of a low- carbon steel, cold rolled 60% and annealed at 540 °C (1005 °F). Note the incubation time is shortened as the penultimate grain size before cold rolling is decreased. (Ref 8) References cited in this section 5. S.R. Goodman and H. Hu, Trans. Met. Soc. AIME, Vol 233, 1965, p 103; Vol 236, 1966, p 710 6. B. Bay and N. Hansen, Met. Trans. A, Vol 10, 1979, p 279; Vol 15A, 1984, p 287 7. A.R. Jones and N. Hansen, in Recrystallization and Grain Growth of Multiphase and Particle Containing Materials, N. Hansen, A.R. Jones, and T. Leffers, Ed., Riso National Laboratory, Denmark, 1980, P 19 8. D.A. Witmer and G. Krauss, Trans. ASM, Vol 62, 1969, p 447 Grain Growth After recrystallization is complete that is, when the polygonized matrix is replaced by the new strain-free grains further annealing increases the average size of the grains. The process, known as grain growth, is accomplished by the migration of grain boundaries. In contrast to recrystallization, the boundary moves toward its center of curvature. Some of the grains grow, but others shrink and vanish. Because the volume of the specimen is a constant, the number of the grains decreases as a consequence of grain growth. The driving force for grain growth is the grain-boundary free-energy, which is substantially smaller in magnitude than the driving energy for recrystallization. According to the growth behavior of the grains, grain growth can be further classified into two types: normal or continuous grain growth and abnormal or discontinuous grain growth. The latter has also been termed exaggerated grain growth, coarsening, or secondary recrystallization. Normal or continuous grain growth occurs in pure metals and single-phase alloys. During isothermal growth, the increase in the average grain diameter obeys the empirical growth law, which can be expressed as D = Kt n where D ¨is the average grain diameter, t is the annealing time, and K and n are parameters that depend on material and temperature. Therefore, when D and t are plotted on a logarithmic scale, a straight line should be obtained, with K as the intercept and n the slope. The value of n, the time exponent in isothermal grain growth, is usually less than, or at most equal to, 0.5. A typical example for isothermal grain growth in zone-refined iron is shown in Fig. 25. The deviation from a straight-line relationship for very short annealing times at low temperatures is due to recrystallization, and that for long annealing times at high temperatures is due to the limiting effect of the sheet specimen thickness. Fig. 25 Normal grain growth in zone- refined iron during isothermal anneals. Closed circles represent specimens for which statistical analysis of grain-size and grain-shape distributions was conducted. One of the structural characteristics during normal grain growth is that the grain size and grain-shape distributions are essentially invariant; that is, during normal grain growth, the average grain size increases, but the size and shape distributions of the grains remain essentially the same before and after the growth, differing only by a scale factor. Figures 26 and 27 show, respectively, the size and shape distributions of the grains in zone-refined iron after normal grain growth at 650 °C (923 K) for various lengths of time. The data points fit the same distribution curves. Therefore, to a first approximation, normal grain growth is equivalent to photographic enlargement. Fig. 26 Grain-size distribution in zone- refined iron during isothermal grain growth at 650 °C (923 K), using a scalar-adjusted grain diameter for each specimen. The plot indicates that the grain- size distribution remains essentially unchanged during normal grain growth. Fig. 27 Grain-shape distribution in zone- refined iron during isothermal grain growth at 650 °C (923 K), using the number of sides of individual grains. The plot indicates that the grain- shape distribution remains essentially unchanged during normal grain growth. During the normal grain growth, the change in texture is small and gradual. Assuming the initial grains are nearly random-oriented, after extensive normal grain growth some weak preferred orientations may be developed among the final grains, depending on such factors as the energies of the free surfaces of the grains. If the initial grains are strongly textured, normal grain growth may be inhibited as a consequence of low mobility of the matrix-grain boundaries (see the next section of this article). Figure 28 shows the grain aggregate of a zone-refined iron specimen after normal grain growth at 800 °C (1470 °F) for 12 min. The size and shape distributions of these grains are essentially the same as those of the much finer grains before growth. Fig. 28 Zone-refined iron, cold rolled to a moderate reduction and annealed for recrystallization for several cycles to refine the penultimate grain size without introducing preferred orientation. Micrograph shows grain structure after normal grain growth at 800 °C (1470 °F) for 12 min. 2% nital. 45× [...]... at 2 0-2 5° from ND toward RD(e) Mg, Co [0001] P ND(e) Ti-Al (>2% Al) (0001) Be, Hf, Zr, Ti, Ti-Nb, Ti-Ta, Ti-Zr [0001] at 2 0-4 0° from ND toward TD, P RD(e) ortho U, 500 °C (932 °F) {1 4 6} + {103} Recrystallization after cold rolling fcc Al, Au, Cu, Cu-Ni, Fe-Cu-Ni, Ni, Ni-Fe, Th Ag, Ag-30Au, Ag-1Zn, Cu-( 5-3 9Zn), Cu-( 1-5 Sn), Cu-0.5... 69,000 MPa (10 × 106 psi) c-axis 5° from tensile axis 10 MPa (1450 psi) c-axis 45° from tensile axis 1.7 MPa (247 psi) a-axis 33 × 1 0-6 /K (59.4 × 1 0-6 /°F) b-axis -6 .5 × 1 0-6 /K (-1 1.7 × 1 0-6 /°F) c-axis 6 × 1 0-4 Ωm (60, 000 μΩ· cm) 90° from c-axis 15 × 1 0-4 Ωm (150, 000 μΩ· cm) c-axis 1.8 T (8, 000 G)(b) Yield strength Thermal expansion Electrical resistivity Magnetic-flux density (a) Magnesium Uranium... Ga oC8 A12 α-Mn cI58 A13 β-Mn cP20 A15 W3O cP8 A20 α-U oC4 B1 NaCl cF8 B2 CsCl cP2 B3 ZnS cF8 B4 ZnS hP4 B81 AsNi hP4 B82 InNi2 hP6 B9 HgS hP6 B10 PbO tP4 B11 γ-CuTi tP4 B13 α-NiS hR6 B16 GeS oP8 B17 PtS tP4 B18 CuS hP12 B19 β'-AuCd oP4 B20 FeSi cP8 B27 BFe oP8 B31 MnP oP8 B32 NaTl cF16 B34 PdS tP16 B35 CoSn hP6 B37 SeTl tI16 Be CdSb oP16 Bf (B33) ζ -CrB oC8 Bg BMo tI16 Bh WC hP2 Bi γ'-CMo (AsTi) hP8... percentage of component: Al, Pb, Co-10Fe, Cu-8Al (< 10%); Au (15%); Ni, 4Mo-79Ni-17Fe, Cu, Cu-2Al, Cu-4Al, Co35Ni(2 5-3 5%); Co-40Ni (50%); Ag (>90%) (d) Binary Cu alloys containing more than 4% Al, 3.5% As, 5% Ge, 0.5% Mg, 4% Mn, 1% P, 3% Sb, 3% Sn, or 10% Zn (e) ND, normal direction; RD, rolling direction; TD, transverse direction Fig 10 Actual {111} and {200} pole figures for electrolytic tough... lines in micrographs, and are two-dimensional defects of lower energy than large-angle grain boundaries Twin boundaries, therefore, are less effective as sources, and sinks, of other defects and are less active in deformation and corrosion than are ordinary grain boundaries Textbooks and reference books, such as Ref 6, 7, and 8 list the indices of twinning planes (shear planes) and the directions of... (Bravais) lattices and their Hermann-Mauguin and Pearson symbols System Space lattice HermannMauguin symbol Pearson symbol Triclinic (anorthic) Primitive P aP Monoclinic Primitive P mP Base-centered(a) C mC Primitive P oP Base-centered(a) C oC Face-centered F oF Body-centered I oI Primitive P tP Body-centered I tI Hexagonal Primitive R(b) hP Rhombohedral Primitive R hR Cubic Primitive P cP Face-centered F... Ni-Fe, Th Ag, Ag-30Au, Ag-1Zn, Cu-( 5-3 9Zn), Cu-( 1-5 Sn), Cu-0.5 Be, Cu-0.5Cd, Cu-0.05P, Co-10Fe {100} {113} bcc Mo Same as deformation texture Fe, Fe-Si, V {111}< 2 11>, and {001} + {112} with < 1 10> 15° from RD(e) Fe-Si {110} after two-step rolling and annealing (Goss method); also {110}, {100} after high-temperature anneal (>1100 °C, or 2012 °F) Ta {111}< 2 11> W, + {112} Ag, Yb, Ni-15Mo, Ni-50Co, Co-10Fe, 1 8-8 stainless steel, Cu alloys(d) {110} < 1 12> + spread around {110} . Uranium b-axis -6 .5 × 10 -6 /K (-1 1.7 × 10 -6 /°F) c-axis 6 × 10 -4 Ωm (60, 000 μΩ· cm) Electrical resistivity Tellurium 90° from c-axis 15 × 10 -4 Ωm (150, 000 μΩ· cm) Magnetic-flux density. (microband region or transition bands) between the main deformation bands. See also Fig. 16. 5% nital. 400× Fig. 14 Fe- 3Si single crystal, cold rolled 80% in the(001)[100] orientation and annealed. inset, indicate similarity in feature between these grain-boundary bands and the transition bands described earlier. These grain-boundary bands obviously would not form at every grain boundary,