VYSOKÁ ŠKOLA BÁŇSKÁ TECHNICKÁ UNIVERZITA OSTRAVA FAKULTA METALURGIE A MATERIÁLOVÉHO INŽENÝRSTVÍ. Thermo-mechanical forming processes.

Transcription

1 VYSOKÁ ŠKOLA BÁŇSKÁ TECHNICKÁ UNIVERZITA OSTRAVA FAKULTA METALURGIE A MATERIÁLOVÉHO INŽENÝRSTVÍ Thermo-mechanical forming processes Study Support doc. Ing. Radim Kocich, Ph.D. Ostrava 2015

2 Title: Thermo-mechanical forming processes Code: Author: doc. Ing. Radim Kocich, Ph.D. Edition: first, 2015 Number of pages:58 Academic materials for the Progressive technical materials study programme at the Faculty of Metallurgy and Materials Engineering. Proofreading has not been performed. Execution: VŠB - Technical University of Ostrava 2

3 1. Thermo-mechanical forming processes (controlled forming) Thermo-mechanical forming controls all the manufacturing conditions with the aim to achieve the required state of structure. It primarily includes composition of steel and thermo-mechanical and time parameters of forming and possible cooling. Suitable structural states enable to achieve higher mechanical properties while maintaining favourable brittleness, fatigue characteristics and weldability. From the structural point of view, the most significant factors are grain and sub-grain size, content of the individual structural components, amount and dispersity of precipitates, dislocations density and structure of solid solution. The basic presupposition for achieving high strength properties is an increased dislocations density; movement of dislocations during an external loading is slowed by internal barriers. The character of these barriers and their distribution is especially important for toughness properties, since these enable to relax peak tension stresses by dislocations movement before the stresses exceed the strength of the material. Barriers with a low permeability include high-angle grain boundaries, non-plastic structural particles and non-coherent matrices including Lomer-Cottrel barriers. The opposite properties have low- or mid-angle grain boundaries, plastic particles and coherent matrices. Time to study: 20 hours Aims After study of this chapter you will be able to define current methods of thermo-mechanical forming and its basic mechanisms and division describe low and high temperature thermo-mechanical processing, isoforming, deformation aging, forming at low temperatures, SHT process, direct rolling, rolling with hot batch 3

4 describe basic types of controlled rolling, define controlled cooling determine mechanical properties of control-formed material define basic physical metallurgical processes in controlled rolling Lecture 1.1. Methods of thermo-mechanical processing of steel The amount of steel processed thermo-mechanically per year is increasing. It can be assumed that a percentage increase of only 10% (approx. 500kt/year) in production with the use of thermo-mechanical methods would create a financial profit of approx. 200 million CZK for the Czech Republic. (Short note: Modern production lines technically enable application of direct thermo-mechanical processing. However, longterm conceptual directed research would be needed so that modern literature-based knowledge, experience from conferences and plastometric and semi-operational (laboratory) experimental results would be fully applicable.) Other advantages of thermo-mechanical forming (TF) will be further mentioned also for aluminium and other non-ferrous metals. In the future, it seems essential to expand these methods of processing. There are several factors that promote such methods. First is the increase in capability and decrease in cost of computers of all kinds, which means simulations can be performed on standard PCs. Second factor is ever deeper understanding of specific phenomena of TF which enables not only a quantitative description of a process, but also to simulate forming operations on common computers. The third factor is a gradual application of TF principles (and also processes which lead to a controlled microstructure development) in an ever increasing number of commercially made products. A directed algorithm is finally becoming a successful process for highly specific products and is expanding its application to more common materials. As has been stated earlier, TF enables to achieve a specific pre-defined microstructure that imparts specific mechanical and physical properties. This path is different from 4

5 traditional forming processes. TF requires control and interaction of at least these mechanisms (among others): dislocation yield and creep recrystallization grain growth phase transformations precipitation particle thickening Most of these structural changes appear dynamically, i.e. during the course of deformation or statically, i.e. after deformation. Each of the above mentioned parameters can be studied individually even if the basic rules are not entirely clear. Furthermore, during TF, synthesis of these basic parameters occurs as well as their mutual interaction. Especially in the area of large deformations, relatively high temperatures and high deformation speeds, these phenomena truly are a great unknown so far. TF can be divided into several individual methods. One of the options could be the following: Grain refinement for increase of strength and toughness without any other further thermal processing. Grain refinement for achieving superplastic behaviour. Texture control (minimization of peaks, improved ductility, magnetic properties ). Inside each of these categories are processes, which must be solved concurrently. It requires laboratory simulation as well as computer modelling. For TMP, two types of steel are usually used: steels with phase changes (including martensitic), steels with precipitation of precipitates from oversaturated solid solution 5

6 (dispersion strengthening) Current TMP methods a) forming before austenite phase change low- and high-temperature TMP b) forming during austenite phase change isoforming c) forming after austenite phase change deformation aging, low-temperature forming Low Temperature TMP (LTMP) LTMP is based on austenitization with rapid cooling into the range of high stability of metastable austenite, forming with high degree of deformation ( over 0.6) and quenching into martensite. Quenching is usually followed by tempering at a lower temperature. In steels for LTMP, sufficient stability of overcooled austenite is required. This makes chrome and molybdenum alloyed steels suitable candidates. The main advantage of LTMP is an increased strength at sufficient plasticity. High-temperature TMP (HTMP) The principle is austenitization, forming just over Ar3 with subsequent quenching and tempering at lower temperatures. It is important that large deformations of austenite and subsequent quenching occur without recrystallization of austenite. For HTMP, steels with slow progress of recrystallization are suitable. In comparison with LTMP, HTMP is technologically simpler, deformation can be performed in the temperature range between to 800 C. HTMP does not allow such increase in strength as LTMP, but is does increase plasticity and fatigue resistance. Isoforming (TBP) Is based on austenitization and cooling down to the temperature of the perlite nose in the IRA diagram. Steel is deformed during the entire transformation and subsequently cooled on air. Tempering is not required. The resulting structure is characterized by formation of fine sub-grains of ferrite and fine globular carbides. Deformation temperatures are usually somewhere between C, this is sufficient for the 6

7 process of polygonization. The deformation is over 0.6. TBP does not significantly increase the strength properties. However, its main advantage is an increase in toughness and decrease in transit temperatures. Deformation aging (DA) It is processing of either conventionally treated or LTMP processed steel, which is then formed at a temperature at which aging occurs. For example, a processing with = 0.02 deformation and temperature between 150 and 200 C. It increases strength, but at the same time plasticity and impact strength are decreased. Low temperature forming (LTF) The initial state is either a state after quenching or a state after annealing. Deformations after quenching help to achieve strength of up to 3130 MPa. The aforementioned TMP methods are mostly used for the common types of steel. These methods do not sufficiently utilize control of structure development through slowed recrystallization and precipitation kinetics. Nevertheless, newly developed and introduced methods of controlled rolling of micro-alloyed steels are utilized to make ferrite grains finer through a transformation from deformed austenite. Slowed recrystallization of austenite is achieved through the combined effect of micro-alloying elements in a solid solution and the interaction of controlled precipitation Basic types of controlled rolling (CR) rolling in the lower region of austenite (in two or three stages) rolling in the region of two-phase structure method a) complemented by controlled cooling controlled rolling with quenching and tempering special methods of CR - for example the SHT (Sumitomo High Toughness) method 7

8 A schematic development of structures during controlled rolling methods can be seen in figure 1. Figure 1: A schematic of microstructure development during rolling procedures Properties and characteristics of structure during thermomechanical processing The main factors that influence yield strength and transition temperature are: strengthening of solid solution through substitution and interstitial elements size of ferrite grain precipitates in ferrite, content of perlite and its inter-lamellar distance Peierls-Nabarro stress ( 40 MPa) During the transformation of from recrystallized austenite, it is possible to express the yield strength K using the Hall-Petch equation (1). 1/ 2 K 0 k. d (1) where 0 is internal stress, d is grain size, ky is coefficient expressing the theoretical value of barrier effect of grain boundaries against movement of dislocations [MPa.mm l/2 ]. 8

9 The value of 0 consists of several components and for the transformation of from non-recrystallized austenite can be expressed by the following relation: 0 S IN Z PR PER PN (2) If the deformation is ended in the + dual-phase area, the final structure consists of recrystallized ferrite grains and subgrains. Since all the grain boundaries in such a structure restrict the movement, the d value has to be replaced with the de effective grain size value (Eq. 3). d 1 e d 1 r 1. f d.1 f r c r (3) where de -1 is length of boundaries for a unit length, dr -1 is length of high-angle boundaries, dc -1 is length of low-angle boundaries, fr is volume fraction of recrystallized ferrite Misorientation of grain boundary angle, types of grain boundaries One of the first steps during TMP is to modify a coarse as-cast structure. One of the aspects according to which structure grain boundaries can be distinguished is the character of their rotation and/or tilt (also called misorientation). Several basic types of grain boundaries can be observed (Figure 2). Figure 2: Possible types of grain boundaries and their energies. A tilted boundary with a low misorientation angle consists of a field of parallel edge dislocations. A higher amount of dislocations causes higher misorientation angle and a consequent increase in energy. Nevertheless, the speed of the increase decreases at higher energy or amount of dislocations due to mutual annihilation of dislocation fields. Besides the misorientation angle and distance between dislocations, the dependency of 9

10 grain boundary energy on misorientation angle θ can also be disturbed by dislocations cores overlapping. This then creates a transfer between low-angle grain boundaries (LAGBs) and high-angle grain boundaries (HAGBs). Special boundaries are locations in which two crystals overlap with a relatively small curvature of inter-atomic layers. This phenomenon is also the basis of the CSL (coincident site lattice) model, which in principle describes overlapping of selected locations in lattices. For example, if every 5 th lattice point is overlapped, then the boundary is denoted as Σ5 etc. Except several exceptions the predicted limits for cubic lattices are Σ29 or Σ33. Nevertheless, the CLS index only rarely depicts the exact boundary. The CSL model is also valid only for 2D systems, while in real conditions the material system is 3D. Another fact is that the data were derived for cubic systems, whereas for non-cubic systems the situation is more complicated Deformation bands Plastic deformation is influenced by changes of the stress state imparted by interactions of dislocations with grain boundaries and obstacles. Actually, each individual grain within the material deforms in an entirely heterogenous way influenced by the neighboring grains being deformed. This was proven by microhardness measurements resulting into different values for individual grains. The final deformation substructure is dependent not only on the temperature and way of deformation, but also on grains orientations and local deformations. This means that the tensors of deformation can differ inside one single grain due to accommodation of the grain to different properties of its neighborhoods (Figure 3). With progressing deformation, especially during deformation under cold conditions, heterogeneity of slip in grains can increase due to misorientations, which tend to develop between individual deformation bands. By this reason the grains subsequently divide into domains (cell blocks) and deformation bands. 10

11 Figure 3: Various types of deformation inhomogeneities; a) slip bands, b) deformation bands, c) dislocation boundaries, d) dislocations. However, the term deformation bands is also sometimes used for coarse bands evident in some big grains caused by divergences of lattice rotations. Neighboring bands tilt away from each other in different directions and become subsequently divided by zones of micrometers sizes with high orientation gradients called transition bands. Cell bands are divided by walls, which increase misorientation only slowly. These can be called as dislocation boundaries, geometrically necessary boundaries (GNBs) or rotational walls. Dislocation cells and dislocation walls are in principle transformation characteristics which continuously appear and disappear during deformation due to random interactions of dislocations. Such boundaries are denoted as random (incidental) dislocation boundaries. On the other hand, walls of cell blocks and transition bands are permanent characteristics of a deformed substructure resulting from plastic deformation Shear bands Bands of localized shear can develop in strongly deformed metals within several scales, from macroscopic, i.e. in the range of centimeters, to microscopic, i.e. in the range of micrometers. When a decrease in hardening rate during plastic flow occurs, plastic deformation starts to localize into shear bands developing from originally wide diffusion bands. Due to progressing deformation these bands rapidly densify into localized zones of intensive shear, which consequently leads to failure. Before the occurrence of 11

12 macroscopic bands shear bands can occur within a grain and subsequently develop through several grains. Alloys hardened with fine precipitates, fine twins or alloys with high dislocation density are especially prone to development of this type of shear bands Structure development Significantly different conditions for structure development of steel occur with different deformation temperatures. Region I recrystallization The original grain size is a function of temperature and material metallurgical character. Refinement of austenite grains is given by the particular deformation-recrystallization cycle. Nuclei for the transformation occur only on boundaries of austenite grains. The final size of ferrite grains is limited by a certain limit value and its further decrease beyond this limit using the deformation-recrystallization cycle within region I is not possible. Eventually, relatively coarse ferrite grains develop after transformation. Region II without recrystallization Forming within region I results into achievement of a limit grain size. Further grain refinement can be achieved by forming in region II. Austenite grains elongate as a consequence of restricted recrystallization and bands with higher dislocation density and higher inner energy and instability develop inside them. Ferrite nuclei develop not only on austenite grain boundaries, but also within deformation bands deformation bands have an effect similar to the effect of grain boundaries during the transformation. Region III austenite ferrite Further grain refinement can be achieved by forming in region III (dual-phase). Ferrite grains deformed after the occurrence of transformation cannot recrystallize anymore and subgrains occur during subsequent recovery. A decreased solubility of Mn and V in ferrite after the transformation fastens stress-induced precipitation, which contributes to pile-up of dislocations and sub-boundaries. Grain growth is also suppressed in the dual-phase region. The result of forming in region III is therefore a mixed structure consisting of equiaxed ferrite grains, grains not deformed after the 12

13 transformation, grains with a lower dislocation density (soft) and subgrains with a higher dislocation density (hard) Structural features during controlled forming Dissolution of carbides and nitrides, grain growth One of the aims of heating to a forming temperature is to dissolve the largest possible amount of precipitates into a solid solution. At the same time, conditions for critical grain growth must not be reached. Solubility of precipitates (nitrides and carbides) can be determined using corresponding equations and can also be influenced by other elements present in steels, especially Mn and Si. Si decreases solubility of Ti in steels. Titanium bonded in stable compounds does not influence solubility of carbides. It is therefore necessary to consider only free Ti. Grain growth is influenced by the amount of impurities and alloying elements present in solid solution and by second phase particles. These reduce grain areas and pin grain boundaries. Increasing amount of Nb and Ti in steels increases critical temperature for grain growth. Vanadium does not restrict grain growth so significantly as Ti and Nb. Vanadium carbides are less stable and dissolve at lower temperatures Strain hardening during TMP Strain hardening during controlled forming is controlled by the amount of energy imposed into a material as a result of plastic deformation. It is a consequence of an increase in stress caused by dislocation movement more or less restricted by obstacles (precipitates, grain boundaries etc.), which they have to surpass during their movement. Stacking fault energy of a deformed material defines the possibility of dislocations to be dissociated into partial dislocations and stacking faults within slip planes. A low stacking fault energy (lower than 50 mj/m 2 ) usually results into dissociation of full dislocations into wide partial dislocations. This subsequently limits their movement only to slip planes and alternative dislocation movements, such as localized cross-slip, out of these planes become extremely difficult (i.e. improbable). At lower strains, these dislocations usually accumulate in slip planes, interact with dislocations in different slip systems and 13

14 form planar dislocation structures, which consequently lead to a development of high localized stress gradients. At higher strains, materials with a low stacking fault energy usually tend to form microtwins (or possibly localized martensitic transformations) and shear bands. They are therefore prone to form evident planar faults of crystal lattice or local orientations. Depending on the stacking fault energy it is possible to divide the development of deformation stress into the four following stages. Stage 1: Movement of dislocations within their slip planes is usually restricted and no mutual interactions occur. By this reason, the velocity of strain hardening is very low. Plastic deformation results into rotation of crystals, which leads to their reorientation and multiple slip, which results into a stronger interaction of dislocations in stage 2. Stage 1 is neglectable for commercial polycrystalline materials since the movement of the first dislocations is limited by grain boundaries, at which dislocations usually pile-up. The resistance to overcome these boundaries can be expressed using a modified Hall-Petch relation (4) ( D ) k D (4) where is deformation resistance for a very large grain, coefficient k1 is usually 0.7 for carbon steel, 0.2 to 0.4 for HCP lattice metals, 0.07 to 0.1 for Cu and Al, respectively. Presence of carbon in a solid solution significantly increases the k coefficient value. Stage 2: Mutual interactions of dislocations in various slip systems cause fast multiplication of dislocations and a subsequent high and approximately constant hardening rate. Development of 3D fields consisting of dislocation multipoles (Taylor networks) occur in metals with low stacking fault energies, while dislocation tangles develop and often form into cellular patterns in metals with high stacking fault energies. Typical strain values characterizing this stage are Stage 3: At strain values approximately up to 1 the deformation curve becomes parabolic. Strain hardening rate decreases progressively down to values approximately an order of 14

15 magnitude lower than at stage 2. In this stage, the effect of the dislocation multiplication process is reduced by dislocation annihilation (dynamic recovery due to localized crossslip, climb or mutual elimination of segments of opposite signs). Microstructure development is towards a clearly defined cell substructure consisting of dislocation cell walls. However, this results into a significant decrease in dislocation density inside cells. The individual cell walls are at first created by a complex of dislocation tangles, which subsequently reduce their thicknesses with progressing deformation (individual dislocations move closer to each other). Cells dimensions reduce during deformation from several micrometers to tenths of micrometers. At the same time misorientation of individual cells increases from the original value of 1 up to 3 to 4 and even more. Stage 4: At strains higher than 1 many grains disintegrate into bands of various orientations divided by transition zones and grain boundaries. A formation of lamellar structure consisting of disoriented microbands parallel to the rolling direction occurs at very high strains. Strain hardening rate in this stage is not very high, although it remains almost constant. Therefore, increase in deformation stress can occur at high strains. Crystal defects, especially dislocations generated by plastic deformation, have high elastic energies, which are stored in the deformed material, especially in stress fields surrounding dislocations. Dislocations mobility can among others be influenced by the amount of dissolved atoms. At moderately elevated temperatures enabling diffusion, dissolved atoms can start to move and then pin on slowly moving dislocations, by which very strong interactions become enabled (usually between room temperature and 300 C). Reduced dislocations mobility caused by segregations of dissolved atoms leads to heterogeneous deformation with the occurrence of Lüders bands and Portevein- LeChatelier (PLC) effects (serrations on stress-strain curves) (Figure 4). Except small and medium strains, the deformation is unlikely to be homogenous and usually leads to localization of strain, typically to the form of bands. This is caused by stacking fault energy, deformation temperature, method of deformation and others. Nevertheless, it is necessary to distinguish two groups of deformation inhomogeneities. The first group consists of deformation bands. These are relatively homogenously deformed regions inside a grain characterized by a different slip system than the 15

16 neighboring regions within the grain. The second group is shear bands. These are areas of very strong localized slip within a grain, as well as through many grains. Both the groups lead to dissociation of a gran into subgrains, although by different mechanisms. Figure 4: Stress-strain curve with evident PLC effect (Al-Mg alloy) Precipitation Defects caused by plastic deformation support diffusion of micro-alloying elements and nucleation of precipitates. As a consequence, a quick strain-induced type of precipitation occurs. Moreover, plastic deformation decreases solubility of micro-alloying elements, which also fastens precipitation. The influence of plastic deformation (ε = 0.09) on solubility of Nb(CN) was expressed by Yamamoto (Eq. 5). log (5) Nb. C N 1, / T Strain-induced precipitation occurring during controlled rolling is especially important during softening processes. While coarse non-dissolved particles (>0.1 m) already present in the structure before deformation concentrate strain in their vicinity and therefore support recrystallization, strain-induced precipitates cause its significant delay. Recrystallization is also inhibited by micro-alloying atoms present in solid solution. The differences in size and electron structure of the micro-alloying atoms and the matrix support their segregation in stacking faults areas, which consequently 16

17 changes stacking fault energies. This causes inhibition of dislocation movement and redistribution of dislocations, which makes nucleation of recrystallization nuclei more difficult. The substantial delaying effect of strain-induced precipitates on recrystallization kinetics can be primarily explained by the fact that precipitation occurs especially on sub-boundaries of deformed austenite. The most effective inhibitors of grain boundaries movement during recrystallization are fine strain-induced precipitates with the diameter smaller than 6nm. Development of recrystallization processes is driven by particles coarsening rate during deformation. The region of precipitates formation is usually between 850 and 1000 C. The coarsening rate at the temperature of 1000 C is significant, while at 950 C it decreases substantially and coarsening is not evident even at longer dwelling times at the temperature of 850 C. For an effective inhibition of recrystallization at least 0.02% of Nb in solid solution is necessary. Precipitation kinetics in austenite and ferrite can generally be described by ARA diagrams. The T0 temperature is the temperature at which the Nb(CN) and Fe3C precipitation curves cross each other. Precipitation of Fe3C at temperatures below Ar1 and at the same time above T0 is improbable, since precipitation of Nb(CN) occurs faster due to its lower free energy when compared to Fe3C. At temperatures below T0 the diffusion ability of Nb in ferrite as a control mechanism of precipitation on the phase boundary of γ/α is lower. Therefore, Nb(CN) precipitates non-uniformly throughout the entire material volume. During forming in the region between the Ar3 and Ar1 temperatures precipitates segregate in parallel layers. Origination of these layers is most probably in precipitation on γ/α phase boundaries during transformation. Such precipitates do not significantly increase strength. Incubation time for static precipitation is in the order of magnitude of 10 2 s at 950 C. For strain-induced precipitation this time shortens by one order of magnitude, while it shortens even by two orders of magnitude for dynamic precipitation. 17

18 Recovery Recovery occurs when a non-equilibrium concentration of lattice defects (point and line defects) is decreased, usually by annealing at an appropriate temperature. Point defects are already recovered at relatively low temperatures (below 0.3 Tt), i.e. below the temperatures of TMPs of most materials. Recrystallization dominates at higher temperatures, although recovery also occurs since kinetics is higher at higher temperatures. For metals with higher stacking fault energies (e.g. Al), processes such as dislocation cross-slip and local annihilation proceed easily, which supports recovery. On the other hand, for metals with low stacking fault energies (e.g. austenitic steels, -brass), the recovery process is not much evident before recrystallization. The same phenomenon applies for alloys with a high amount of atoms dissolved in solid solution reducing dislocation mobility. Recovery can occur immediately after plastic deformation, but also during deformation. Nevertheless, recovery does not influence the appearance of microstructure or crystallographic texture. It influences properties, such as hardness, dislocation density, size and misorientation of subgrains. However, the changes can only hardly be detected. Mechanisms applied during recovery - annihilation of point defects (vacancies, interstitions) using diffusion e.g. into dislocations - mutual annihilations of dislocations (applies for closely located dislocations of opposite signs or dipoles, for which a low amount of dislocation climb or crossslip is needed) - rearrangement of free dislocations and random dislocations into dislocation walls or sub-boundaries (polygonization) - coalescence of sub-boundaries walls during subgrains growth It is very difficult to distinguish the process of diffusion of dislocation cells into clearly defined subgrain boundaries from the process of grain growth. This is especially due to the heterogeneity of dislocation structure in deformed polycrystalline materials. The late stage of recovery (formation of clearly defined sub-boundaries) is often the first 18

19 stage of nucleation during recrystallization, which can lead to a fast recrystallization stopping further recovery. Structural changes during recovery Besides the first fast decrease in volume of point defects, the main structural changes can be divided as follows. - rearrangement of dislocations into cell structures (for metals with high stacking fault energies and most of metals formed under hot conditions this proceeds together with deformation) - elimination of free dislocations inside cells - rearrangement of complicated dislocation structure of cell walls into arranged subgrain boundaries mostly by annihilation of redundant dislocations and rearrangement of other dislocations into low-energy configurations (Figure 5) - subgrains growth (during progressing annealing subgrain growth occurs, since this leads to decrease in internal energy) Figure 5: Changes of dislocation structure during recovery from random dislocation tangles through cell substructures to subgrains [1] Extended recovery/continuous recrystallization Recrystallization is a discontinuous process, during which absorption of deformed/recovered areas by non-deformed grains via HAGBs movement occurs. It is possible to suppress discontinuous recrystallization in some cases, which subsequently leads to a relatively fine and uniform grain structure. This process features homogenous subgrains growth and is often called extended recovery. It is not clear whether this is 19

20 caused only by LAGBs movement, or whether HAGBs movement also contributes to this process. Both these processes can occur locally (inside individual grains), as well as globally. - local extended recovery was detected e.g. for Al, low-carbon steels and Zr-based alloys, for grains with specific orientations, which are usually not affected by local deformation, i.e. are subjected to a homogenous deformation - global recovery/continuous recrystallization was detected for dual-phase alloys and for cases of extreme plastic deformations (severe plastic deformations SPD methods) Secondary phase precipitates can pin and stabilize sub-grains structure. A subsequent dissolution/coarsening of precipitates can lead to a homogenous subgrains growth or extended recovery. This phenomenon was originally identified as in situ or continuous recrystallization. SPD processes are often applied to achieve ultra-fine grained (UFG) structure prepared using continuous recrystallization (static or dynamic). For these processes a large deformation together with nucleation/dissolution of secondary phases has a significant effect. The result can be extended recovery/continuous recrystallization, but also geometrically necessary dynamic recrystallization (see below). In this case, recovery can make localized deformation easier Recrystallization Primary recrystallization (also called discontinuous recrystallization) is often defined as a process of nucleation and growth. The driving force for nucleation during recrystallization is far lower than e.g. during solidification. Every large subgrain or a relatively organized region inside a deformed grain can be considered as a possible recrystallization nucleus based only on the driving force or relative differences in the accumulated energies between the possible nucleus and its surroundings. Whether the possible nucleus is real (active) depends on its ability to grow, which depends especially on the presence of growth-facilitating boundaries (i.e. the presence and mobility of HAGBs, the mobility of which is far larger than the mobility of LAGBs). LAGBs created by deformation/recovery occurring usually between subgrains feature a very low mobility, 20

21 whereas boundaries with higher misorientations (10-15 ) feature a very high mobility. This results in a nucleation originating in a rapid growth of a very small minority of subgrains, which subsequently evolve into new growing grains. The first necessary condition is that a subgrain should have (or quickly achieve) a local misorientation larger than 15. A very quick growth of a very small amount of subgrains comparing to a slow growth of the remaining subgrains imparts a heterogeneous character to this casual type of recrystallization, which is then depicted as nucleation and growth. Typically, a region without deformation (i.e. without grain boundaries or misorientations) larger than a certain (critical) size is considered to be a recrystallized grain. To achieve this critical size (usually in the order of magnitude of micrometers), nucleation and restricted or only local growth are necessary. It is very difficult to distinguish whether a possible nucleus is active or not. The difference between an original subgrain (possible nucleus) and a final recrystallized grain for typical metals is times the original size. The probability of finding an active nucleus in a deformed/recrystallized metal matrix is Sources of recrystallized grains - Deformed grains In deformed grains, recrystallized grains with more or less similar orientations as the original grains can develop. Recrystallized grains can however develop also from deformation bands of similar orientations. Such bands are effective recrystallization sources due to substantial localized deformations inside the grains. This non-uniform localization of deformation leads to fragmentations of the bands, as well as to large inhomogeneities in the values of imposed strain. - Shear bands Shear bands passing through several grains can also be sources of recrystallized grains. This is usually related to a high imposed energy and therefore large variations in relative misorientations and a possible presence of growth-facilitating boundaries. Shear bands are typical especially for metals with low stacking fault energies. - Particle stimulated nucleation (PSN) 21

22 Dislocations can be pinned by relatively large particles insusceptible to slip. This pinning and subsequent growth of dislocation density leads to formations of locally deformed regions with large developments of misorientations around secondary phase particles. Consequently, recrystallization occurs in these locally deformed regions due to large differences in the imposed energies. This type of recrystallization is usually depicted as PSN (particle stimulated nucleation). Locally deformed regions around particles are then depicted as PSN grains with random orientations. Low-temperature annealing performed after deformation often results in a highly random orientation, which is caused by the existence of inner and outer locally deformed regions, or by the influence of PSN and deformation bands. One of the mechanisms by which a recrystallized grain with a new orientation can originate is also recrystallization twins. These occur especially in metals with low stacking fault energies (Cu, austenitic steel). As was already mentioned, particles are very important during recrystallization, since they can pin grain boundaries during their movement. However, there are more types of the pinning effect. - Pinning by Zener drag. Low-energy boundaries, such as CSL boundaries, have low Zener drag pinning effect. - Pinning using elements dissolved in solid solution (solute drag). The effect is low for CSL boundaries. - Orientation pinning. Growing grains can be pinned by locally deformed regions of similar orientations, i.e. by a presence of LAGBs. Possible types of recrystallization - Dynamic recrystallization Under certain conditions, a structure can recrystallize during deformation, i.e. dynamically. Occurrence of this type of recrystallization is theoretically possible also during deformation under cold conditions, although it is practically observed only rarely (for very pure metals). Dynamic recrystallization can either be discontinuous, geometrical or based on a progressive rotation of subgrains. The latter two types are 22

23 based on the presupposition of imposed strain with a limited or restricted movement of HAGBs. - Discontinuous dynamic recrystallization Figure 6 depicts typical stress-strain curves describing deformation under cold and hot conditions. For a deformation under hot conditions, the shape of the curve can be influenced by the strain hardening rate depending on dynamic recovery or recrystallization (i.e. discontinuous dynamic recrystallization). The dynamic recovery curve is typical for metals with high stacking fault energies (e.g. Al, low-carbon steel), where a deformation resistance steady state occurs after the first stage of strain hardening. Evolutions of microstructures during dynamic recovery and dynamic recrystallization are schematically depicted in Figure 7. During dynamic recovery the original grains deform, while sub-boundaries stay more or less equiaxed. The structure is thus dynamic and adapts continuously to the increasing strain. The recovery process is significantly slower for metals with low stacking fault energies (e.g. austenitic steel, Cu), which can enable accumulation of imposed energy. At a certain critical strain, dynamically recrystallized grains start to occur on the boundaries of the original grains. As a result, a chain structure occurs at grain boundaries. With an increasing strain more nuclei get activated and new recrystallized grains occur. At the same time, already recrystallized grains are deformed again. After a certain amount of imposed strain, the stress-strain curve gets saturated. The microstructure then consists of a dynamic mixture of grains with various dislocation densities. 23

24 Figure 6. Typical deformation curves describing deformation under cold and hot conditions. Figure 7. Evolution of microstructures during deformation under hot conditions: (a) recovery, (b) dynamic recrystallization and discontinuous dynamic recrystallization Subgrain size (originating from dynamic recovery), as well as grain size (originating from dynamic recrystallization) increase with temperature and decrease with increasing strain rate. Contrary to dynamic recovery, dynamic recrystallization also includes metadynamic recrystallization. During this type of recrystallization, a recrystallized nucleus develops or grows dynamically under hot conditions. Nevertheless, only growth occurs during subsequent static annealing. - Geometrical dynamic recrystallization Grains with serrated boundaries developed as a result of dynamic recovery can mutually pin when their sizes achieve their thicknesses. This process can create a microstructure appearing as dynamically recrystallized. However, it is a result of deformation and recovery. - Dynamic recrystallization via progressive subgrains rotation During deformation, subgrains in the vicinities of already existing boundaries can be subjected to more substantial rotations comparing to central grains areas. Development of HAGBs can occur at high strains. Although the exact mechanism is not clear, it is supposed that this is due to a combination of inhomogeneous plastic deformation, 24

25 accelerated recovery (in regions in the vicinities of grain boundaries) and grain boundary sliding Grain coarsening (growth) Grain growth after recrystallization is controlled especially by surface energy or grain boundary energy. This driving force is by two orders of magnitude lower than driving force for recrystallization. A more correct description for this grain growth process is grain coarsening (sometimes also secondary recrystallization). There are two basic types of grain coarsening: normal and abnormal. During normal coarsening, the main mechanism is elimination of finest grains, while the grain size distribution remains almost constant. During abnormal coarsening, growth of several grains inside a pinned structure occurs. The pinning effect is usually caused by particles of other phases or by low misorientation angles (i.e. low grain boundaries mobility). The grain growth driving force is substantially lowered if the grain size achieves approximately ½ of the thickness of the deformed sample. This is due to a reduced grain boundary radius. At the grain boundaries equilibrium state, the force has to act perpendicular to the sample surface. Moreover, grooves develop in the intersection of grain boundaries with the surface, which restrict grain boundaries movement. Due to a very high number of influencing factors it is not possible to determine in advance when and how abnormal grain growth will occur Alternative deformation mechanisms Although slip is the dominant deformation mechanism for most of metal materials, activation of different mechanisms is possible under certain conditions. Among these are twinning, creep, grain boundary sliding (GBS), and deformation related to phase transformations. Various deformation mechanisms can be active during deformation under hot conditions depending on the temperature and applied stress. To define areas, in which the individual deformation mechanisms can be expected to occur, deformation mechanisms maps are used. The maps practically bring information about the influence 25

26 of the temperature and applied stress on the strain rate and dominant deformation mechanism. An example of such a map is shown in Figure 8. Figure 8: Deformation mechanisms map for W (grain size 10 m). [1] The upper boundary of the diagram is created by theoretical shear stress. This is stress that should be applied in a case of perfect deformation of crystals with no defects by mutual collective translation of crystal planes. Deformation proceeds via dislocations movement in their slip planes, which is supported by high temperature, i.e. by the process of dynamic recovery caused by cross-slip and climb. In the lower part of the map, creep region can be observed. This region can further be divided into various sub-regions, each of which is characterized by a creep submechanism. Creep can be caused by dislocation movement or diffusion. Diffusion creep has two types, Nabarro and Coble creeps. Twinning can also be mentioned among deformation mechanisms occurring under these conditions. However, it is typical only for several materials - twinning, as well as GBS, are not included in Figure 8. The reason for neglecting GBS is since GBS in polycrystalline materials is usually enabled by diffusion creep. Therefore GBS is already included inside the deformation maps diffusion creep areas. GBS also occurs during superplastic forming since it is supported by very fine grain size and elevated temperatures. A concrete example can be tungsten, which deforms by diffusion creep and the strain rate of 10-9 s -1 at the temperature of 1560 C and a relatively low stress of 10 MPa. This 26

27 means that a deformation of 10% would take 3 years, which is very important information for certain applications. On the other hand, at the stress of 100 MPa, the material will deform by dislocation creep and the strain rate of 10-6 s -1. At the stress of 1000 MPa, the casual dislocation slip will dominate and the strain rate will be s -1. There is a presupposition that high-temperature dislocation creep is not different from mechanisms occurring during conventional deformation under hot conditions Creep Creep can be characterized as a deformation response of a material to long-time loading. It comprises continuous elongation terminated by fracture. Creep is typical for high temperatures, it occurs at low strain rates and relatively low stresses. Creep deformation curve can be divided into three stages: Stage 1 primary creep When a load is applied, a material is deformed and new dislocations are generated. The dislocations then make obstacles in movement for each other, which causes decrease in creep rate (Eq. 6) 0 d / dt (6) Stage 2 secondary creep (steady-state creep, creep at constant rate) In this area, strain hardening rate and dynamic recovery rate are equal. A constant load is applied, which results into a constant strain rate after a certain time period. Stage 3 tertiary creep Voids develop in the material, which causes increase in effective stress and also creep rate. Voids growth quickly results into material fracture. At low temperatures, strain hardening is dominant and the steady-state stage is usually not reached. At high temperatures or high stresses, development of first voids can occur even during the first stage, i.e. steady-state stage is either very short, or completely missing. Creep mechanisms 1. Diffusion creep 27

28 During loading of a material (plastic deformation) the material tries to deform via movement of vacations through a crystal along its boundaries. Therefore the deformation proceeds by a creation of vacancies on these boundaries due to tensile stresses. At the same time, destruction of vacancies on boundaries by the influence of compression stresses occurs. This can be expressed by Eq D b 14 D (7) 2 v ktd D Dv where is atom volume, D is grain size, Db, Dv are diffusion coefficients for volume and grain boundaries diffusion, respectively, is grain boundaries area effective for grain boundaries diffusion. This equation assumes that ancillary processes of nucleation and annihilation at grain boundaries and GBS require a similar negligible energy and stress. However, if this condition is not met and volume diffusion dominates, Nabarro or Nabarro-Herring creep occur. Contrariwise, Coble creep occurs in cases of dominant diffusion at grain boundaries. 2. Dislocation creep If a stress higher than corresponding to the diffusion creep region in a deformation map is applied, another deformation mechanism is activated. Although deformation is realized via dislocation slip and climb, diffusion is a process controlling dislocations recovery by climbing. Depending on the temperature and stress, four main groups of dislocation creep mechanism can be distinguished. 1 slip and climb controlled by volume diffusion 2 slip and climb controlled by pipe diffusion 3 Harper-Dorn creep 4 power law breakdown The first two mechanisms can be characterized using the power law relation (Eq. 8). n DbGb 0 A (8) kt G For higher stresses, the following energetic dissolution relation is valid (Eq. 9). 28

29 0 K exp (9) A linear dependence between stress and strain rate - Harper-Dorn creep - applies for some materials Grain boundary sliding and superplasticity Superplasticity is an ability of crystalline materials to extremely plastically deform when subjected to tensile stress. The elongation is usually on the order of several hundred percent. Superplasticity in most cases occurs at high temperatures and low strain rates, while the necessary stress is rather low (to 20 MPa). Although this phenomenon has already been widely studied, its application is still limited due to the requirement of very low strain rates. Applicability of superplastic forming is therefore especially for fabrication of components with complex shapes and with limited production capacities, which can be for example specific components for aerospace industry. Several types of superplasticity can be distinguished. One of them is superplasticity of fine-grain structures, since fine grains are among the necessary conditions for this type (GBS is the main mechanism). Besides, there are transformation superplasticity, superplasticity induced by inner stresses and superplasticity at high strain rates (hyperelasticity). During GBS, mutual sliding of two grains occurs due to shear stress. In contrast to dislocation slip during which grains get elongated in the direction of deformation, grains more or less maintain their dimensions during GBS. For superplasticity forming GBS is the entirely dominant mechanism. During GBS, cavities form in the regions where grains slide along each other. It is therefore clear that during superplastic forming a mechanism providing time suppressing (delaying) of cavities development has to be active. To enable continuation of sliding, the material has to be shifted (usually by diffusion creep) from regions where grains overlap (regions subjected to compression stress) to regions with tendencies to create voids (regions subjected to tensile stress). 29

30 Conditions of superplasticity The first condition of superplastic forming for ultra-fine grained materials is the condition of high m (strain rate sensitivity) coefficient value, which is usually defined by Eq. 10 for a concrete deformation temperature. 0 log / log (T, ) (10) High values of m minimalize formations of necks during tensile tests (i.e. localization of deformation), expected in the absence of strain hardening. The only possible mechanism of strain hardening during diffusion creep GBS is grain coarsening. The second condition is an increase in temperature to a sufficient value. The critical value is usually a temperature higher than half of the melting temperature of the particular material. The third condition is the already mentioned strain rate (usually to 10-2 s -1, preferably lower). Mathematically it can be described by Eq m k( ) (11) The presupposition is therefore the existence of two possible deformation mechanisms: dislocation slip and GBS enabled by diffusion creep. The m coefficient is reciprocal to the n coefficient. Since m is higher for GBS than for dislocation sliding (m>1/n), the curve for GBS has steeper slope than for dislocation creep (Figure 9). At higher strain rates, the required stress is lower for dislocation sliding and therefore this mechanism dominates. Contrariwise, GBS dominates for lower strain rates. In the region of dominant GBS enabled by diffusion creep, m is a function of strain rate. Figure 9: GBS vs. dislocation sliding for m DS < m GBS. [1] 30