ALL THIS MICROSTRUCTURES ARE TAKEN FROM DIFFERENT WEBSITES (ONLY FOR EDUCATION PERPOSE)
Tuesday, September 11, 2007
Microstructures of Different Steels
ALL THIS MICROSTRUCTURES ARE TAKEN FROM DIFFERENT WEBSITES (ONLY FOR EDUCATION PERPOSE)
Saturday, September 8, 2007
Temper Rolling,Luders Bands
when the stress strain curves of metals deformed in ension are plotted, two basic types of curves are observed.One curve exhibit sharp yield point,where the stress rises with almost negliible plastic deformation toa point calld upper yield point.At this point the material begin to yield , with a simultaneous drop in flow stressrequired for contnued deformation .This new yield point is called lower yield point and corresponds to an appreciable plastic strain at an almost constant stress.Evenually the metal starts to harden with in incrase in the stress necessary for additional deformation. After this occurs , there is little difference between the appearance of the stress strain curves for metals with a yield point and those without it.
The sharp yield point is an especially important effect because it occurs in iron and in low carbon steels . Its existance is an considerable concern to manufacturers who stamp or draw thin sheets of these material in forming such objects as automobile bodies.Th significance of the yield point is this :once plastic deformation starts in a given area ,the metal at this point is effectively softened and suffers a relatively large plastic deformation . This deformation then spreads into the material adjoining the region which has yielded because of the stress concentration at the boundary between the deformed and undeformed areas.In general,deformation starts at positions of stress concentration as discrete bands of deformed material ,called LUDERS BAND.
The effect of using metals containing a yield point is to develop a roughened surface.This surface results from the uneven spread of the LUDERS BANDS which leaves striation on the surface ,commonly called STRECHER STRAINS.
Annealed steel sheet is often given a slight reduction in thickness by rolling which amounts to about a 1% strain.This is called TEMPER ROLLING(SKIN PASS), and it produces a very large number of luders band neclei in the sheet . When the metals is deformed latter into a finished product , these small bands grow,but because of there small size and close proximity to each other ,the resulting surface roughening is greatly reduced.
Rferences:i)Physical Metallurgy,R.E REEDHILL,
ii)http://www.ccr.buffalo.edu/etomica/app/modules/sites/MaterialFracture/Images/SSPicture3.jpg
Tuesday, September 4, 2007
Tempered Martensite
Introduction
Tempering is a term historically associated with the heat treatment of martensite in steels. It describes how the microstructure and mechanical properties change as the metastable sample is held isothermally at a temperature where austenite cannot form. The changes during the tempering of martensite can be categorised into stages. During the first stage, excess carbon in solid solution segregates to defects or forms clusters within the solid solution. It then precipitates, either as cementite in low-carbon steels, or as transition iron-carbides in high-carbon alloys. The carbon concentration that remains in solid solution may be quite large if the precipitate is a transition carbide. Further annealing leads to stage 2, in which almost all of the excess carbon is precipitated, and the carbides all convert into more stable cementite. Any retained austenite may decompose during this stage. Continued tempering then leads to the coarsening of carbides, extensive recovery of the dislocation structure, and finally to the recrystallisation of the ferrite plates into equiaxed grains.
This is a useful description but it is revealing to consider first, the factors responsible for driving the process in the first place.
Deviation from Equilibrium
Tempering is a process in which the microstructure approaches equilibrium under the influence of thermal activation. It follows that the tendency to temper depends on how far the starting microstructure deviates from equilibrium. It is interesting therefore to consider how metastable a material can be, before dealing specifically with martensite. Turnbull characterised metastability in terms of the unit RTm where R is the universal gas constant and Tm is the absolute melting temperature. This coarse unit is a measure of the thermal energy in the system at the melting temperature; it represents a large amount of energy, typically in excess of 20,000 J mol-1.
Table 1: Degree of metastability Metastable structure RTm
Supersaturated solutions 1
Amorphous metal 0.5
Modulated films, nanostructures 0.1
Supersaturated solutions are prominent in this list and the extent of metastability depends both on the excess concentration and on the equilibrium solubility. It can be demonstrated that excess carbon which is forced into solution in martensite is the major contributor to the stored energy of martensite.
The calculations presented in Table 2 show the components of the stored energy of martensite in a typical low--alloy martensitic steel Fe-0.2C-1.5Mn wt%. It is necessary to define a reference state, which is here taken to be an equilibrium mixture of ferrite, graphite and cementite, with a zero stored energy. Graphite does not in fact form because it is too slow to precipitate; the effect of replacing the graphite with cementite is to increase the stored energy by some 70 J mol-1.
When transformations occur at low temperatures, it is often the case that substitutional elements like manganese and iron cannot diffuse during the time scale of the experiment, whereas carbon is still mobile. The transformation then happens in such a way that the Fe/Mn ratio is maintained constant whilst the carbon redistributes subject to this constrain, until its chemical potential becomes uniform. This is known as paraequilibrium. Unlike the equilibrium state, because the iron and manganese atoms are trapped during transformation, their chemical potentials are no longer uniform. This adds a further 315 J mol-1 to the stored energy.
When bainite forms, the transformation mechanism is displacive, there is a shape deformation, which leads to an additional 400 J mol-1 of stored energy. Since there is no diffusion during transformations, but the carbon partitions following growth, the total stored energy is that for the paraequilibrium state added to the strain energy term, giving a net value of 785 J mol-1.
Martensite is not only a diffusionless transformation, but it frequently occurs at low temperatures where its virgin microstructure is preserved. Even the carbon remains trapped in the product crystal. Furthermore, the strain energy term associated with martensite is greater at about 600 J mol-1 because the plates tend to have a larger aspect ratio (thickness/length). There may also be twin interfaces within the martensite plates, which cost about 100 J mol-1. The trapping of carbon inside the martensite adds a further 629 J mol-1, which makes the total stored energy in excess of 1700 J mol-1!
Table 2: Stored energies of a variety of microstructures
Phase Mixture in Fe-0.2C-1.5Mn wt% at 300 K Stored Energy / J mol-1
1. Ferrite, graphite and cementite 0
2. Ferrite and cementite 70
3. Paraequilibrium ferrite and paraequilibrium cementite 385
4. Bainite and paraequilibrium cementite 385+400=785
5. Martensite 385+600+100+629=1714
6. Mechanically alloyed ODS metal 55
The stored energy becomes even larger as the carbon concentration is increased (Figure 1).
Figure 1: The free energy due to the trapping of carbon in martensite, as a function of its carbon concentration. The results are for a temperature of 473 K.
Virgin Microstructure
The virgin microstructure obtained immediately after quenching from austenite consists of plates or laths of martensite which is supersaturated with carbon. In the vast majority of steels, the martensite contains a substantial density of dislocations which are generated during the imperfect accommodation of the shape change accompanying the transformation. The plates may be separated by thin films of retained austenite, the amount of untransformed austenite becoming larger as the martensite-start temperature MS is reduced.
a) Transmission electron micrograph of as-quenched martensite in a Fe-4Mo-0.2C wt% steel. The mottled contrast within the plates is due to a high density of dislocations. (b) Corresponding dark-field image showing the distribution of retained austenite.
Carbon Atoms
Carbon is an interstitial atom in ferritic iron, primarily occupying the octahedral interstices. There are three such interstices per iron atom. At a typical concentration of 0.4 wt% or about 2 at%, less than 1% of these interstices are occupied by carbon. Furthermore, there is a strong repulsion between carbon atoms in nearest neighbour sites. This means that carbon atoms almost always have an adjacent interstitial site vacant, leading to a very high diffusion coefficient when compared with the diffusion of substitutional solutes. In the latter case, the substitutional vacancy concentration is only 10-6 at temperatures close to melting, and many orders of magnitude less at the sort of temperatures where martensite is tempered. It follows that carbon diffuses much faster than substitutional atoms (including iron), as illustrated below.
a) A carbon atom in an octahedral interstice in body-centered cubic iron
b) The ratio of the diffusivity of a substitutional atom to that of carbon in body-centered cubic iron.
Given that carbon is able to migrate in martensite even at ambient temperature, it is likely that some of it redistributes, for example by migrating to defects, or by rearranging in the lattice such that the overall free energy is minimised.
Precipitation of Iron Carbides
In high-carbon steels, the precipitation of excess carbon begins with the formation of a transition carbide, such as ε (Fe2.4C). ε-carbide can grow at temperatures as low as 50oC. Indeed, most of the iron carbides can precipitate at low temperatures, well below those associated with the motion of substitutional solutes. This is because they grow by a displacive mechanism which does not require the redistribution of substitutional atoms (including iron); carbon naturally has to partition. This corresponds to a process known as paraequilibrium transformation in which the iron to substitutional solute ratio is maintained constant but subject to that constraint, the carbon achieves a uniform chemical potential.
Martensite is said to be supersaturated with carbon when the concentration exceeds its equilibrium solubility with respect to another phase. However, the equilibrium solubility depends on the phase. The solubility will be larger when the martensite is in equilibrium with a metastable phase such as ε carbide. Some 0.25 wt% of carbon is said to remain in solution after the precipitation of ε-carbide is completed.
Although most textbooks will begin a discussion of tempering with this first stage of tempering, involving the redistribution of carbon and precipitation of transition carbides, cementite can precipitate directly. This is particularly the case when the defect density is large. Trapped carbon atoms will not precipitate as transition carbides but cementite is more stable than trapped carbon
a) Transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 190oC for 1 hour. The carbon has in this case precipitated as fine particles of cementite. (b) Corresponding dark-field image showing the distribution of retained austenite, which has not been affected by the temperin
Decomposition of Retained Austenite
Tempering at higher temperatures, in the range 200-300oC for 1 h induces the retained austenite to decompose into a mixture of cementite and ferrite. When the austenite is present as a film, the cementite also precipitates as a continuous array of particles which have the appearance of a film.
Dark field transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 295oC for 1 hour. Only the cementite is illuminated. The film of cementite at the martensite plate boundaries is due to the decomposition of retained austenite.
Further Tempering
Tempering at even higher temperatures leads to a coarsening of the cementite particles, with those located at the plate boundaries growing at the expense of the intra-plate particles. The dislocation structure tends to recover, the extent depending on the chemical composition. The recovery is less marked in steels containing alloying elements such as molybdenum and chromium.
Bright field transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 420oC for 1 hour.
The original article is in the following link
www.msm.cam.ac.uk/phase-trans/2004/Tempered.Martensite/tempered.martensite.html
Tempering is a term historically associated with the heat treatment of martensite in steels. It describes how the microstructure and mechanical properties change as the metastable sample is held isothermally at a temperature where austenite cannot form. The changes during the tempering of martensite can be categorised into stages. During the first stage, excess carbon in solid solution segregates to defects or forms clusters within the solid solution. It then precipitates, either as cementite in low-carbon steels, or as transition iron-carbides in high-carbon alloys. The carbon concentration that remains in solid solution may be quite large if the precipitate is a transition carbide. Further annealing leads to stage 2, in which almost all of the excess carbon is precipitated, and the carbides all convert into more stable cementite. Any retained austenite may decompose during this stage. Continued tempering then leads to the coarsening of carbides, extensive recovery of the dislocation structure, and finally to the recrystallisation of the ferrite plates into equiaxed grains.
This is a useful description but it is revealing to consider first, the factors responsible for driving the process in the first place.
Deviation from Equilibrium
Tempering is a process in which the microstructure approaches equilibrium under the influence of thermal activation. It follows that the tendency to temper depends on how far the starting microstructure deviates from equilibrium. It is interesting therefore to consider how metastable a material can be, before dealing specifically with martensite. Turnbull characterised metastability in terms of the unit RTm where R is the universal gas constant and Tm is the absolute melting temperature. This coarse unit is a measure of the thermal energy in the system at the melting temperature; it represents a large amount of energy, typically in excess of 20,000 J mol-1.
Table 1: Degree of metastability Metastable structure RTm
Supersaturated solutions 1
Amorphous metal 0.5
Modulated films, nanostructures 0.1
Supersaturated solutions are prominent in this list and the extent of metastability depends both on the excess concentration and on the equilibrium solubility. It can be demonstrated that excess carbon which is forced into solution in martensite is the major contributor to the stored energy of martensite.
The calculations presented in Table 2 show the components of the stored energy of martensite in a typical low--alloy martensitic steel Fe-0.2C-1.5Mn wt%. It is necessary to define a reference state, which is here taken to be an equilibrium mixture of ferrite, graphite and cementite, with a zero stored energy. Graphite does not in fact form because it is too slow to precipitate; the effect of replacing the graphite with cementite is to increase the stored energy by some 70 J mol-1.
When transformations occur at low temperatures, it is often the case that substitutional elements like manganese and iron cannot diffuse during the time scale of the experiment, whereas carbon is still mobile. The transformation then happens in such a way that the Fe/Mn ratio is maintained constant whilst the carbon redistributes subject to this constrain, until its chemical potential becomes uniform. This is known as paraequilibrium. Unlike the equilibrium state, because the iron and manganese atoms are trapped during transformation, their chemical potentials are no longer uniform. This adds a further 315 J mol-1 to the stored energy.
When bainite forms, the transformation mechanism is displacive, there is a shape deformation, which leads to an additional 400 J mol-1 of stored energy. Since there is no diffusion during transformations, but the carbon partitions following growth, the total stored energy is that for the paraequilibrium state added to the strain energy term, giving a net value of 785 J mol-1.
Martensite is not only a diffusionless transformation, but it frequently occurs at low temperatures where its virgin microstructure is preserved. Even the carbon remains trapped in the product crystal. Furthermore, the strain energy term associated with martensite is greater at about 600 J mol-1 because the plates tend to have a larger aspect ratio (thickness/length). There may also be twin interfaces within the martensite plates, which cost about 100 J mol-1. The trapping of carbon inside the martensite adds a further 629 J mol-1, which makes the total stored energy in excess of 1700 J mol-1!
Table 2: Stored energies of a variety of microstructures
Phase Mixture in Fe-0.2C-1.5Mn wt% at 300 K Stored Energy / J mol-1
1. Ferrite, graphite and cementite 0
2. Ferrite and cementite 70
3. Paraequilibrium ferrite and paraequilibrium cementite 385
4. Bainite and paraequilibrium cementite 385+400=785
5. Martensite 385+600+100+629=1714
6. Mechanically alloyed ODS metal 55
The stored energy becomes even larger as the carbon concentration is increased (Figure 1).
Figure 1: The free energy due to the trapping of carbon in martensite, as a function of its carbon concentration. The results are for a temperature of 473 K.
Virgin Microstructure
The virgin microstructure obtained immediately after quenching from austenite consists of plates or laths of martensite which is supersaturated with carbon. In the vast majority of steels, the martensite contains a substantial density of dislocations which are generated during the imperfect accommodation of the shape change accompanying the transformation. The plates may be separated by thin films of retained austenite, the amount of untransformed austenite becoming larger as the martensite-start temperature MS is reduced.
a) Transmission electron micrograph of as-quenched martensite in a Fe-4Mo-0.2C wt% steel. The mottled contrast within the plates is due to a high density of dislocations. (b) Corresponding dark-field image showing the distribution of retained austenite.
Carbon Atoms
Carbon is an interstitial atom in ferritic iron, primarily occupying the octahedral interstices. There are three such interstices per iron atom. At a typical concentration of 0.4 wt% or about 2 at%, less than 1% of these interstices are occupied by carbon. Furthermore, there is a strong repulsion between carbon atoms in nearest neighbour sites. This means that carbon atoms almost always have an adjacent interstitial site vacant, leading to a very high diffusion coefficient when compared with the diffusion of substitutional solutes. In the latter case, the substitutional vacancy concentration is only 10-6 at temperatures close to melting, and many orders of magnitude less at the sort of temperatures where martensite is tempered. It follows that carbon diffuses much faster than substitutional atoms (including iron), as illustrated below.
a) A carbon atom in an octahedral interstice in body-centered cubic iron
b) The ratio of the diffusivity of a substitutional atom to that of carbon in body-centered cubic iron.
Given that carbon is able to migrate in martensite even at ambient temperature, it is likely that some of it redistributes, for example by migrating to defects, or by rearranging in the lattice such that the overall free energy is minimised.
Precipitation of Iron Carbides
In high-carbon steels, the precipitation of excess carbon begins with the formation of a transition carbide, such as ε (Fe2.4C). ε-carbide can grow at temperatures as low as 50oC. Indeed, most of the iron carbides can precipitate at low temperatures, well below those associated with the motion of substitutional solutes. This is because they grow by a displacive mechanism which does not require the redistribution of substitutional atoms (including iron); carbon naturally has to partition. This corresponds to a process known as paraequilibrium transformation in which the iron to substitutional solute ratio is maintained constant but subject to that constraint, the carbon achieves a uniform chemical potential.
Martensite is said to be supersaturated with carbon when the concentration exceeds its equilibrium solubility with respect to another phase. However, the equilibrium solubility depends on the phase. The solubility will be larger when the martensite is in equilibrium with a metastable phase such as ε carbide. Some 0.25 wt% of carbon is said to remain in solution after the precipitation of ε-carbide is completed.
Although most textbooks will begin a discussion of tempering with this first stage of tempering, involving the redistribution of carbon and precipitation of transition carbides, cementite can precipitate directly. This is particularly the case when the defect density is large. Trapped carbon atoms will not precipitate as transition carbides but cementite is more stable than trapped carbon
a) Transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 190oC for 1 hour. The carbon has in this case precipitated as fine particles of cementite. (b) Corresponding dark-field image showing the distribution of retained austenite, which has not been affected by the temperin
Decomposition of Retained Austenite
Tempering at higher temperatures, in the range 200-300oC for 1 h induces the retained austenite to decompose into a mixture of cementite and ferrite. When the austenite is present as a film, the cementite also precipitates as a continuous array of particles which have the appearance of a film.
Dark field transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 295oC for 1 hour. Only the cementite is illuminated. The film of cementite at the martensite plate boundaries is due to the decomposition of retained austenite.
Further Tempering
Tempering at even higher temperatures leads to a coarsening of the cementite particles, with those located at the plate boundaries growing at the expense of the intra-plate particles. The dislocation structure tends to recover, the extent depending on the chemical composition. The recovery is less marked in steels containing alloying elements such as molybdenum and chromium.
Bright field transmission electron micrograph of martensite in a Fe-4Mo-0.2C wt% steel after tempering at 420oC for 1 hour.
The original article is in the following link
www.msm.cam.ac.uk/phase-trans/2004/Tempered.Martensite/tempered.martensite.html
annealing of steel part 2
Introduction ANNEALING is a generic term denoting a treatment that consists of heating to and holding at a suitable temperature followed by cooling at an appropriate rate, primarily for the softening of metallic materials. Generally, in plain carbon steels, annealing produces a ferrite-pearlite microstructure (Fig. 1). Steels may be annealed to facilitate cold working or machining, to improve mechanical or electrical properties, or to promote dimensional stability. The choice of an annealing treatment that will provide an adequate combination of such properties at minimum expense often involves a compromise. Terms used to denote specific types of annealing applied to steels are descriptive of the method used, the equipment used, or the condition of the material after treatment.
Metallurgical Principles The iron-carbon binary phase diagram (Fig. 2) can be used to better understand annealing processes. Although no annealing process ever achieves true equilibrium conditions, it can closely parallel these conditions. In defining the various types of annealing, the transformation temperatures or critical temperatures are usually used. (See the article "Principles of Heat Treating of Steels" in this
Critical Temperatures. The critical temperatures that must be considered in discussing annealing of steel are those that define the onset and completion of the transformation to or from austenite. For a given steel, the critical temperatures depend on whether the steel is being heated or cooled. Critical temperatures for the start and completion of the transformation to austenite during heating are denoted, respectively, by Ac1 and Ac3 for hypoeutectoid steels and by Ac1 and Accm for hypereutectoid steels. These temperatures are higher than the corresponding critical temperatures for the start and completion of the transformation from austenite during cooling, which are denoted, respectively, by Ar3 and Ar1 for hypoeutectoid steels and by Arcm and Ar1 for hypereutectoid steels. (The "c" and "r" in the symbols are derived from the French words chauffage for heating and refroidissement for cooling.) These critical temperatures converge to the equilibrium values Ae1, Ae3, and Aecm as the rates of heating or cooling become infinitely slow. The positions of the Ae1, Ae3, and Aecm lines are close to the more general (that is, near equilibrium) A1, A3, and Acm lines on the iron-carbon binary phase diagram shown in Fig. 2. Various alloying elements markedly affect these critical temperatures. For example, chromium raises the eutectoid temperature, A1, and manganese lowers it. It is possible to calculate upper and lower critical temperatures using the actual chemical composition of the steel. The following equations will give an approximate critical temperature for a hypoeutectoid steel (Ref 1):
Annealing Cycles
In practice, specific thermal cycles of an almost infinite variety are used to achieve the various goals of annealing. These cycles fall into several broad categories that can be classified according to the temperature to which the steel is heated and the method of cooling used. The maximum temperature may be below the lower critical temperature, A1 (subcritical annealing); above A1 but below the upper critical temperature, A3 in hypoeutectoid steels, or Acm in hypereutectoid steels (intercritical annealing); or above A3 (full annealing).
Because some austenite is present at temperatures above A1, cooling practice through transformation is a crucial factor in achieving desired microstructure and properties. Accordingly, steels heated above A1 are subjected either to slow continuous cooling or to isothermal treatment at some temperature below A1 at which transformation to the desired microstructure can occur in a reasonable amount of time. Under certain conditions, two or more such cycles may be combined or used in succession to achieve the desired results. The success of any annealing operation depends on the proper choice and control of the thermal cycle, based on the metallurgical principles discussed in the following sections.
Subcritical Annealing Subcritical annealing does not involve formation of austenite. The prior condition of the steel is modified by such thermally activated processes as recovery, recrystallization, grain growth, and agglomeration of carbides. The prior history of the steel is, therefore, an important factor. In as-rolled or forged hypoeutectoid steels containing ferrite and pearlie, subcritical annealing can adjust the hardnesses of both constituents, but excessively long times at temperature may be required for substantial softening. The subcritical treatment is most effective when applied to hardened or cold-worked steels, which recrystallize readily to form new ferrite grains. The rate of softening increases rapidly as the annealing temperature approaches A1. Cooling practice from the subcritical annealing temperature has very little effect on the established microstructure and resultant properties. A more detailed discussion of the metallurgical processes involved in subcritical annealing is provided in Ref 2.
lntercritical Annealing
Austenite begins to form when the temperature of the steel exceeds A1. The solubility of carbon increases abruptly (nearly 1%) near the A1 temperature. In hypoeutectoid steels, the equilibrium structure in the intercritical range between A1 and A3 consists of ferrite and austenite, and above A3 the structure becomes completely austenitic. However, the equilibrium mixture of ferrite and austenite is not achieved instantaneously. For example, the rate of solution for a typical eutectoid steel is shown in Fig. 3. Undissolved carbides may persist, especially if the austenitizing time is short or the temperature is near A1, causing the austenite to be inhomogeneous. In hypereutectoid steels, carbide and austenite coexist in the intercritical range between A1 and Acm; and the homogeneity of the austenite depends on time and temperature.
The degree of homogeneity in the structure at the austenitizing temperature is an important consideration in the development of annealed structures and properties. The more homogeneous structures developed at higher austenitizing temperatures tend to promote lamellar carbide structures on cooling, whereas lower austenitizing temperatures in the intercritical range result in less homogeneous austenite, which promotes formation of spheroidal carbides.
Fig. 3 Austenitizing rate-temperature curves for commercial plain carbon eutectoid steel. Prior treatment was normalizing from 875 °C (1610 °F); initial structure, fine pearlite. First curve at left shows beginning of disappearance of pearlite; second curve, final disappearance of pearlite; third curve, final disappearance of carbide; fourth curve, final disappearance of carbon concentration gradients.
Austenite formed when steel is heated above the A1 temperature transforms back to ferrite and carbide when the steel is slowly cooled below A1. The rate of austenite decomposition and the tendency of the carbide structure to be either lamellar or spheroidal depend largely on the temperature of transformation. If the austenite transforms just below A1, it will decompose slowly. The product then may contain relatively coarse spheroidal carbides or coarse lamellar pearlite, depending on the composition of the steel and the austenitizing temperature. This product tends to be very soft. However, the low rate of transformation at temperatures just below A1 necessitates long holding times in isothermal treatments, or very slow cooling rates in continuous cooling, if maximum softness is desired. Isothermal treatments are more efficient than slow continuous cooling in terms of achieving desired structures and softness in the minimum amount of time. Sometimes, however, the available equipment or the mass of the steel part being annealed may make slow continuous cooling the only feasible alternative.
As the transformation temperature decreases, austenite generally decomposes more rapidly, and the transformation product is harder, more lamellar, and less coarse than the product formed just below A1. At still lower transformation temperatures, the product becomes a much harder mixture of ferrite and carbide, and the time necessary for complete isothermal transformation may again increase.
Temperature-time plots showing the progress of austenite transformation under isothermal (IT) or continuous transformation (CT) conditions for many steels have been widely published (Ref 3, 4) and illustrate the principles just discussed. These IT or CT diagrams may be helpful in design of annealing treatments for specific grades of steel, but their usefulness is limited because most published diagrams represent transformation from a fully austenitized, relatively homogeneous condition, which is not always desirable or obtainable in annealing.
In the continuous annealing process, which is discussed in detail in the following article in this Section, an intercritical annealing practice is used to develop dual-phase and tri-phase microstructures. In this practice, the steel is rapidly cooled from the intercritical temperature. The rapid cooling results in the transformation of the pools of austenite to martensite. The final microstructure consists of islands of martensite in a ferritic matrix. Depending upon the alloy content of the austenite pools and the cooling conditions, the austenite may not fully transform and the microstructure will consist of martensite/retained austenite regions in a ferritic matrix.
Cooling after Full Transformation. After the austenite has been completely transformed, little else of metallurgical consequence can occur during cooling to room temperature. Extremely slow cooling may cause some agglomeration of carbides, and consequently, some slight further softening of the steel, but in this regard such slow cooling is less effective than high-temperature transformation. Therefore, there is no metallurgical reason for slow cooling after transformation has been completed, and the steel may be cooled from the transformation temperature as rapidly as feasible in order to minimize the total time required for the operation.
If transformation by slow continuous cooling has been used, the temperature at which controlled cooling may be stopped depends on the transformation characteristics of the steel. However, the mass of the steel or the need to avoid oxidation are practical considerations that may require retarded cooling to be continued below the temperature at which the austenite transformation ceases.
Effect of Prior Structure. The finer and more evenly distributed the carbides in the prior structure, the faster the rate at which austenite formed above A1 will approach complete homogeneity. The prior structure, therefore, can affect the response to annealing. When spheroidal carbides are desired in the annealed structure, preheating at temperatures just below A1 sometimes is used to agglomerate the prior carbides in order to increase their resistance to solution in the austenite on subsequent heating. The presence of undissolved carbides or concentration gradients in the austenite promotes formation of a spheroidal, rather than lamellar, structure when the austenite is transformed. Preheating to enhance spheroidization is applicable mainly to hypoeutectoid steels but also is useful for some hypereutectoid low-alloy steels.
Supercritical or Full Annealing
A common annealing practice is to heat hypoeutectoid steels above the upper critical temperature (A3) to attain full austenitization. The process is called full annealing. In hypoeutectoid steels (under 0.77% C), supercritical annealing (that is, above the A3 temperature) takes place in the austenite region (the steel is fully austenitic at the annealing temperature). However, in hypereutectoid steels (above 0.77% C), the annealing takes place above the A1 temperature, which is the dual-phase austenite-cementite region. Figure 4 shows the annealing temperature range for full annealing superimposed in the iron-carbon binary phase diagram from Fig. 2. In general, an annealing temperature 50 °C (90 °F) above the A3 for hypoeutectic steels and A1 for hypereutectoid steels is adequate.
Austenitizing Time and Dead-Soft Steel. Hypereutectoid steels can be made extremely soft by holding for long periods of time at the austenitizing temperature. Although the time at the austenitizing temperature may have only a small effect on actual hardnesses (such as a change from 241 to 229 HB), its effect on machinability or cold-forming properties may be appreciable. Long-term austenitizing is effective in hypereutectoid steels because it produces agglomeration of residual carbides in the austenite. Coarser carbides promote a softer final product. In lower-carbon steels, carbides are unstable at temperatures above A1 and tend to dissolve in the austenite, although the dissolution may be slow. Steels that have approximately eutectoid carbon contents generally form a lamellar transformation product if austenitized for very long periods of time. Long-term holding at a temperature just above the A1 temperature may be as effective in dissolving carbides and dissipating carbon-concentration gradients as is short-term holding at a higher temperature.
Guidelines for Annealing The metallurgical principles discussed above have been incorporated by Payson (Ref 6) into the following seven rules, which may be used as guidelines for development of successful and efficient annealing schedules:
· Rule 1: The more homogeneous the structure of the as-austenitized steel, the more completely lamellar will be the structure of the annealed steel. Conversely, the more heterogeneous the structure of the asaustenitized steel, the more nearly spheroidal will be the annealed carbide structure
· Rule 2: The softest condition in the steel is usually developed by austenitizing at a temperature less than 55 °C (100 °F) above A1 and transforming at a temperature (usually) less than 55 °C (100 °F) below A1
· Rule 3: Because very long times may be required for complete transformation at temperatures less than 55 °C (100 °F) below A1, allow most of the transformation to take place at the higher temperature, where a soft product is formed, and finish the transformation at a lower temperature, where the time required for completion of transformation is short
· Rule 4: After the steel has been austenitized, cool to the transformation temperature as rapidly as feasible in order to minimize the total duration of the annealing operation
· Rule 5: After the steel has been completely transformed, at a temperature that produces the desired microstructure and hardness, cool to room temperature as rapidly as feasible to decrease further the total time of annealing
· Rule 6: To ensure a minimum of lamellar pearlite in the structures of annealed 0.70 to 0.90% C tool steels and other low-alloy medium-carbon steels, preheat for several hours at a temperature about 28 °C (50 °F) below the lower critical temperature (A1) before austenitizing and transforming as usual
· Rule 7: To obtain minimum hardness in annealed hypereutectoid alloy tool steels, heat at the austenitizing temperature for a long time (about 10 to 15 h), then transform as usual These rules are applied most effectively when the critical temperatures and transformation characteristics of the steel have been established and when transformation by isothermal treatment is feasible.
Annealing Temperatures
From a practical sense, most annealing practices have been established from experience. For many annealing applications, it is sufficient simply to specify that the steel be cooled in the furnace from a designated annealing (austenitizing) temperature. Temperatures and associated Brinell hardnesses for simple annealing of carbon steels are given in Table 2, and similar data for alloy steels are presented in Table 3.
Table 2 Recommended temperatures and cooling cycles for full annealing of small carbon steel forgings
Data are for forgings up to 75 mm (3 in.) in section thickness. Time at temperature usually is a minimum of 1 h for sections up to 25 mm (1 in.) thick; 12 h is added for each additional 25 mm (1 in.) of thickness.
In isothermal annealing to produce a pearlitic structure, particularly in forgings, an austenitizing temperature as much as 70 °C (125 °F) higher than that indicated in Table 4 may be selected in order to decrease the austenitizing time. For most steels, as indicated in Table 4, annealing may be accomplished by heating to the austenitizing temperature and then either cooling in the furnace at a controlled rate or cooling rapidly to, and holding at, a lower temperature for isothermal transformation. Both procedures result in virtually the same hardness; however, considerably less time is required for isothermal transformation.
Uniformity of Temperature. One potential contribution to the failure of an annealing operation is a lack of knowledge of the temperature distribution within the load of steel in the furnace. Furnaces large enough to anneal 18 Mg (20 tons) of steel at a time are not uncommon. In some large forging shops, workpieces can weigh in excess of 270 Mg (300 tons). The larger the furnace, the more difficult it is to establish and maintain uniform temperature conditions throughout the load, and the more difficult it is to change the temperature of the steel during either heating or cooling.
Furnace thermocouples indicate the temperature of the space above, below, or beside the load, but this temperature may differ by 28 °C (50 °F) or more from the temperature of the steel itself, especially when the steel is in a pipe or box, or when bar or strip is packed in a dense charge in a quiescent atmosphere. When these conditions exist, the distribution of temperature throughout the load during heating and cooling should be established by placing thermocouples among the bars, forgings, coils, and so on. A good practice is to spot weld a thermocouple to the workpiece or to use embedded thermocouples (thermocouples placed in holes drilled into the workpiece). Regulation of the furnace during the annealing operation should be based on the temperatures indicated by these thermocouples, which are in actual contact with the steel, rather than on the temperatures indicated by the furnace thermocouples.
Spheroidizing
The majority of all spheroidizing activity is performed for improving the cold formability of steels. It is also performed to improve the machinability of hypereutectoid steels, as well as tool steels. A spheroidized microstructure is desirable for cold forming because it lowers the flow stress of the material. The flow stress is determined by the proportion and distribution of ferrite and carbides. The strength of the ferrite depends on its grain size and the rate of cooling. Whether the carbides are present as lamellae in pearlite or spheroids radically affects the formability of steel. Steels may be spheroidized, that is, heated and cooled to produce a structure of globular carbides in a ferritic matrix. Figure 5 shows 1040 steel in the fully spheroidized condition. Spheroidization can take place by the following methods:
· Prolonged holding at a temperature just below Ae1
· Heating and cooling alternately between temperatures that are just above Ac1 and just below Ar1
· Heating to a temperature just above Ac1, and then either cooling very slowly in the furnace or holding at a temperature just below Ar1
· Cooling at a suitable rate from the minimum temperature at which all carbide is dissolved to prevent reformation of a carbide network, and then reheating in accordance with the first or second methods above (applicable to hypereutectoid steel containing a carbide network) It should be noted that it is difficult to establish consistent designations for critical temperatures. In discussions about heating with prolonged holding, the critical temperatures of interest should be the equilibrium temperatures Ae1 and Ae3. Terminology becomes more arbitrary in discussions of heating and cooling at unspecified rates and for unspecified holding times.
For full spheroidizing, austenitizing temperatures either slightly above the Ac1 temperature or about midway between Ac1 and Ac3 are used. If a temperature slightly above Ac1 is to be used, good loading characteristics and accurate temperature controls are required for proper results; otherwise, it is conceivable that Ac1 may not be reached and that austenitization may not occur.
Low-carbon steels are seldom spheroidized for machining, because in the spheroidized condition they are excessively soft and "gummy," cutting with long, tough chips. When low-carbon steels are spheroidized, it is generally to permit severe deformation. For example, when 1020 steel tubing is being produced by cold drawing in two or three passes, a spheroidized structure will be obtained if the material is annealed for 1
2 to 1 h at 690 °C (1275 °F) after each pass. The final product will have a hardness of about 163 HB Tubing in this condition will be able to withstand severe deformation during subsequent cold forming.
As with many other types of heat treatment, hardness after spheroidizing depends on carbon and alloy content. Increasing the carbon or alloy content, or both, results in an increase in the as-spheroidized hardness, which generally ranges from 163 to 212 HB (Table 4).
Process Annealing
As the hardness of steel increases during cold working, ductility decreases and additional cold reduction becomes so difficult that the material must be annealed to restore its ductility. Such annealing between processing steps is referred to as in-process or simply process annealing. It may consist of any appropriate treatment. In most instances, however, a subcritical treatment is adequate and least costly, and the term "process annealing" without further qualification usually refers to an in-process subcritical anneal. Figure 9 shows the range of temperatures typically used for process annealing. It is often necessary to specify process annealing for parts that are cold formed by stamping, heading, or extrusion. Hotworked high-carbon and alloy steels also are process annealed to prevent them from cracking and to soften them for shearing, turning, or straightening.
Process annealing usually consists of heating to a temperature below Ae1, soaking for an appropriate time and then cooling, usually in air. In most instances, heating to a temperature between 10 and 20 °C (20 and 40 °F) below Ae1 produces the best combination of microstructure hardness, and mechanical properties. Temperature controls are necessary only to prevent heating the material above Ae1 and thus defeating the purpose of annealing. When process annealing is performed merely to soften a material for such operations as cold sawing and cold shearing, temperatures well below Ae1 normally are used and close controls are unnecessary.
In the wire industry, process annealing is used as an intermediate treatment between the drawing of wire to a size slightly larger than the desired finished size and the drawing of a light reduction to the finished size. Wire thus made is known as annealed in process wire. Process annealing is used also in the production of wire sufficiently soft for severe upsetting and to permit drawing the smaller sizes of low-carbon and medium-carbon steel wire that cannot be drawn to the desired small size directly from the hot-rolled rod. Process annealing is more satisfactory than spheroidize annealing for a material that, because of its composition or size (or both), cannot be drawn to finished size because it either lacks ductility or does not meet physical requirements. Also, material that is cold sheared during processing is process annealed to raise the ductility of the sheared surface to a level suitable for further processing.
Annealed Structures for Machining
Different combinations of microstructure and hardness, considered together, are significant in terms of machinability. For instance, Fig. 10 shows that a partially spheroidized 5160 steel shaft was machined (by turning) with much less tool wear and better surface finish than the same steel in the annealed condition with a pearlitic microstructure and a higher hardness. Based on many observations, optimum microstructure for machining steels of various carbon contents are usually as follows:
The type of machining operation is also a factor. For example, certain gears were made from 5160 steel tubing by the dual operation of machining in automatic screw machines and broaching of cross slots. The screw-machine operations were easiest with thoroughly spheroidized material, but a pearlitic structure was more suitable for broaching. A semispheroidized structure proved to be a satisfactory compromise.
Semispheroidized structures can be achieved by austenitizing at lower temperatures, and sometimes at higher cooling rates, than those used for achieving pearlitic structures. The semispheroidized structure of the 5160 steel tubing mentioned above was obtained by heating to 790 °C (1450 °F) and cooling at 28 °C/h (50 °F/h) to 650 °C (1200 °F). For this steel, austenitizing at a temperature of about 775 °C (1425 °F) results in more spheroidization and less pearlite. Medium-carbon steels are much more difficult to fully spheroidize than are high-carbon steels such as 1095 and 52100. In the absence of excess carbides to nucleate and promote the spheroidizing reaction, it is more difficult to achieve complete freedom from pearlite in practical heat-treating cycles.
At lower carbon levels, structures consisting of coarse pearlite in a ferrite matrix often are found to be the most machinable. In some alloy steels, this type of structure can best be achieved by heating to temperatures well above Ac3 to establish a coarse austenite grain size, then holding below Ar1 to allow coarse, lamellar pearlite to form. This process sometimes is referred to as cycle annealing or lamellar annealing. For example, forged 4620 steel gears were heated rapidly in a five-zone furnace to 980 °C (1800 °F), cooled to 625 to 640 °C (1160 to 1180 °F) in a water-cooled zone, and held at that temperature for 120 to 150 min. The resulting structure--coarse, lamellar pearlite in a ferrite matrix--had a hardness of 140 to 146 HB (Ref .
Types of Furnaces
Furnaces for annealing are of two basic types: batch furnaces and continuous furnaces. Within either of these two types, furnaces can be further classified according to configuration, type of fuel used, method of heat application, and means by which the load is moved through, or supported in, the furnace. Other factors that must be considered in furnace selection are cost, type of annealing cycle, required atmosphere, and physical nature of parts to be annealed. In many cases, however, the annealing cycle used is dictated by the available equipment.
Batch-type furnaces are necessary for large parts such as heavy forgings and often are preferred for small lots of a given part or grade of steel and for the more complex alloy grades requiring long cycles. Specific types of batch furnaces include car-bottom, box, bell, and pit furnaces. Annealing in bell furnaces can produce the greatest degree of spheroidization (up to 100%). However, the spheroidizing cycles in bell furnaces are long and last from 24 to 48 h depending on the grade of material being annealed and the size of the load.
Continuous furnaces such as roller-hearth, rotary-hearth, and pusher types are ideal for isothermal annealing of large quantities of parts of the same grade of steel. These furnaces can be designed with various individual zones, allowing the work to be consecutively brought to temperature, held at temperature, and cooled at the desired rate. Continuous furnaces are not able to give complete spheroidization and should not be used for products that require severe cold forming. For more detailed discussion of the types of furnaces available for annealing, see Ref 9 and the article "Types of Heat- Treating Furnaces" in this Volume.
Furnace Atmospheres
Electric furnaces used with air atmospheres, and gas furnaces used with atmospheres consisting of the products of combustion, cannot be regulated for complete elimination of oxidation of the steel being treated. Only atmospheres independent of the fuel are generally considered satisfactory for clean or bright annealing. Excessive oxidation during annealing usually is prevented by the use of controlled atmospheres in conjunction with a suitable furnace that is designed to exclude air and combustion gases from the heating chamber. The gases and gas mixtures used for controlled atmospheres depend on the metal being treated, the treatment temperature, and the surface requirements of the parts being annealed. The need to eliminate decarburization as well as oxidation is often a significant factor in the selection of annealing atmospheres.
The gas most widely used as a protective atmosphere for annealing is exothermic gas. This gas is inexpensive, the raw materials for making it are readily available, and the results obtained with it are generally excellent. Hydrocarbon gases such as natural gas, propane, butane, and coke-oven gas are commonly burned in an exothermic-gas producer, creating a self-supporting, heat-producing combustion reaction. A commonly used exothermic gas mixture contains 15% H2, 10% CO, 5% CO2, 1% CH4, and 69% N2. This gas is used for bright annealing of cold-rolled low-carbon steel strip. It will decarburize medium-carbon and high-carbon steels, however, because of the carbon dioxide and water vapor it contains. Exothermic gas sometimes is refrigerated to reduce its moisture content, particularly in geographic areas where the temperature of the water used for cooling is high. When decarburization of workpiece surfaces must be prevented, water vapor and carbon dioxide must be completely removed from the gas. Purified exothermic gas, with its carbon dioxide and water vapor removed, has many applications in heat treatment of steel without decarburization.
Purified rich exothermic gas, formed by partial combustion, is used for short-cycle annealing and process annealing of medium- and high-carbon steels of the straight-carbon and alloy types. For long-cycle batch annealing, however, this gas is unsuitable because its high carbon monoxide content results in soot deposits on the work and because of the possibility of surface etching as a result of the relatively long time for which the work is in the critical low-temperature range where gas reactions can occur. In short-cycle annealing these effects are minimal, and the high-CO gas is then desirable because of its high carbon potential. The fairly lean purified gas formed by more complete combustion is used for long-cycle annealing of medium- and high-carbon steels of the straight-carbon and alloy types, and for batch and continuous annealing of low-carbon steel strip for tin plating.
Allowable decarburization on spheroidize annealed blanks or coiled rod can be quite restrictive. As long as the furnace has excellent sealing characteristics, low dewpoint exothermic gas can protect the steel from decarburization. Many commercial heat treaters compensate for the sealing problems of furnaces by using a blend of exothermic and endothermic gases. Depending on the carbon content of the stock that is being processed, the blend can be varied. A great deal of caution has to be exercised when such blends are used because endothermic gas forms an explosive mixture with air as the temperature drops below 760 °C (1400 °F).
Other atmospheres commonly used in annealing include endothermic-base, dissociated ammonia, and vacuum atmospheres. Nitrogen-base atmospheres became popular among heat treaters in the 1980s due in part to rising costs of utilities such as natural gas and water. The nitrogen is blended with small percentages of additives such as methane, propane, propylene, and carbon monoxide. For more complete information, see the article "Furnace Atmospheres" in this Volume.
Annealing of Sheet and Strip
In terms of total tonnage of material processed, annealing of sheet and strip during production of steel-mill products represents the major use of annealing. Because such annealing is done to prepare the material for further processing (such as additional cold rolling or fabrication into parts), and because the temperatures employed are usually below the A1 temperature, the more specific terms subcritical annealing and process annealing are appropriate, although common practice is to use the term annealing without qualification.
In annealing of sheet and strip, two techniques predominate: the batch process and the continuous process. In the batch process (also called box annealing), coils or cut lengths of sheet are placed on an annealing base and covered with containers that are sealed to hold the appropriate atmosphere. A furnace is then placed over the covered steel. A protective atmosphere is introduced within the inner covers to protect the steel from oxidation and is circulated through the coils by use of fans and convector plates. Heating is provided by the outer furnace and may be done either through the use of radiant tubes or by direct firing. The charge is heated to the required temperature and held for a period of time that will result in the desired properties. The outer furnace is then removed, and the coils are allowed to cool under the inner covers. When the temperature has been reduced to the point where oxidation of the steel will not occur, the inner covers are removed and the steel is forwarded for further processing.
In the continuous process, steel coils are uncoiled and drawn through a furnace where they are subjected to the annealing cycle under a protective atmosphere. After the sheet or strip has been cooled and removed from the furnace, further inline processing (such as hot dip galvanizing) may be done, or the steel may be cut into sheets. In general, however, the steel is recoiled and then forwarded as in the batch process. For more detailed information, see the article "Continuous Annealing" that immediately follows in this Section.
In addition to the obvious differences in equipment, the batch process and the continuous process differ considerably in several other ways. Batch annealing may require up to a week because of the large mass of material being treated, whereas continuous annealing is accomplished in about five minutes. Differences are also evident in the temperatures employed, with the batch process generally being conducted at lower temperatures. Because in batch annealing it is difficult to ensure that the temperature is uniform throughout the charge (which may consist of several hundred tons of steel), the continuous process offers the potential of more uniform properties. The short annealing times of the continuous process, however, frequently result in hardness levels slightly higher than those of similar material annealed by the batch process.
Cold-Rolled Plain Carbon Sheet and Strip. The usual method of manufacturing cold-rolled sheet and strip is to produce a hot-rolled coil, pickle it to remove scale (oxide), and cold roll it to the desired final gage. Cold rolling may reduce the thickness of the hot-rolled material in excess of 90%, which increases the hardness and strength of the steel but severely decreases its ductility. If any large amount of subsequent cold working is to be done, the ductility of the steel must be restored.
Annealing of the cold-rolled steel normally is designed to produce a recrystallized ferrite microstructure from the highly elongated, stressed grains resulting from cold work. Figure 11 shows the effect of annealing on the microstructure of a low-carbon cold-rolled sheet steel. The cold-rolled microstructure is shown in Fig. 11(a) in contrast to the partially and fully recrystallized microstructure in Fig. 11(b) and 11(c). During heating of the steel, and in the first segment of the holding portion of the cycle, the first metallurgical process to occur is recovery. During this process, internal strains are relieved (although little change in the microstructure is evident), ductility is moderately increased, and strength is slightly decreased.
As annealing continues, the process of recrystallization occurs, and new, more equiaxed ferrite grains are formed from the elongated grains. During recrystallization, strength decreases rapidly, with a corresponding increase in ductility. Further time at temperature causes some of the newly formed grains to grow at the expense of other grains; this is termed grain growth and results in modest decreases in strength and small (but often significant) increases in ductility. Most plain carbon steels are given an annealing treatment that promotes full recrystallization, but care must be taken to avoid excessive grain growth, which can lead to surface defects (such as orange peel) in formed parts.
The rates at which the metallurgical processes noted above proceed are functions of both the chemical composition and the prior history of the steel being annealed. For example, small amounts of elements such as aluminum, titanium, niobium, vanadium, and molybdenum can decrease the rate at which the steel will recrystallize, making the annealing response sluggish and therefore necessitating either higher temperatures or longer annealing times to produce the same properties. Although the presence of these alloying elements is generally the result of deliberate additions intended to modify the properties of the sheet (as in the case of aluminum, titanium, niobium, and vanadium), some elements may be present as residual elements (molybdenum, for example) in quantities great enough to modify the response to annealing. Conversely, larger amounts of cold work (greater cold reductions) will accelerate the annealing response. Therefore, it is not possible to specify a single annealing cycle that will produce a particular set of mechanical properties in all steels; the chemical composition and the amount of cold work also must be taken into account.
Cold-rolled plain carbon steels are produced to a number of different quality descriptions. Commercial quality (CQ) steel is the most widely produced and is suitable for moderate forming. Drawing quality (DQ) steel is produced to tighter mechanical-property restrictions for use in more severely formed parts. Drawing quality special killed (DQSK) steel is produced to be suitable for the most severe forming applications. Typical properties of these grades may be found in the article "Carbon and Low-Alloy Steel Sheet and Strip," in Properties and Selection: Irons, Steels, and High-Performance Alloys, Volume 1 of ASM Handbook, formerly 10th Edition Metals Handbook. Structural quality (SQ) steel is produced to specified mechanical properties other than those for the above three grades.
Typical annealing cycles for all possible combinations of composition, cold reduction, and grade cannot be listed here. However, typical batch-annealing temperatures range from 620 to 690 °C (1150 to 1270 °F) for the coldest point in the charge. Cycle times vary with the grade desired and the size of the charge, but total times (from the beginning of heating to removal of the steel from the furnace) can be as long as one week. Figure 12 shows a typical heating and cooling cycle for batch-annealing coils of low-carbon cold-rolled steel sheet.
Continuous-annealing cycles are of shorter duration and are conducted at higher temperatures than batch-annealing cycles. In some applications, the annealing temperature may exceed A1. Typical cycles are 40 s at 700 °C (1290 °F) for cold-rolled commercial quality steel and 60 s at 800 °C (1470 °F) for drawing quality special killed sheet. Most continuous annealing of cold-rolled sheet includes an overaging treatment designed to precipitate carbon and nitrogen from solution in the ferrite and to reduce the likelihood of strain aging. Overaging for 3 to 5 min at 300 to 450 °C (570 to 840 °F) accomplishes the desired precipitation of carbon and nitrogen.
Batch annealing and continuous annealing differ slightly in the properties they produce. Typical average properties of batch-annealed and continuous-annealed commercial quality plain carbon steel are as follows:
High-strength cold-rolled sheet and strip are growing in importance due to their high load-bearing capacities. Strength of sheet and strip can be increased through modifications of chemical composition and/or selection of different annealing cycles, but these methods result in decreased ductility. Plain carbon steels, produced by conventional techniques, may be batch annealed or continuous annealed under conditions that result only in recovery or partial recrystallization. Typical batch-annealing cycles of this type employ soak temperatures of 425 to 480 °C (800 to 900 °F) and various soak times. High-strength low-alloy (HSLA) steels containing alloying elements such as niobium, vanadium, and titanium also may be produced as cold-rolled grades. The additional alloying produces a stronger hot-rolled steel, which is strengthened even more by cold rolling. Cold-rolled HSLA steels may be recovery annealed to produce higher-strength grades or recrystallization annealed to produce lower-strength grades. Successful production of cold-rolled HSLA steel requires selection of the appropriate combination of steel composition and hot-rolled strength, amount of cold reduction, and type of annealing cycle. For more information on HSLA steels, see the article "High-Strength Structural and High-Strength Low-Alloy Steels," in Volume 1 of ASM Handbook, formerly 10th Edition Metals Handbook.
Another series of high-strength sheet steels are the dual-phase steels. These steels are generally annealed for a short period (usually less than 5 min) in the intercritical range, followed by rapid cooling. The resulting microstructure is 10 to 20% martensite by volume in a matrix of ferrite. The continuous-annealing process is ideal for producing dual-phase sheet grades (more details are in the next article in this Section). Dual-phase steels are unique in that they deform by a continuous yielding behavior because the martensite is a continuous source of dislocations during plastic deformation (see the article "Dual-Phase Steels" in Volume 1 of ASM Handbook, formerly 10th Edition Metals Handbook). Most other low-carbon steels that display a yield point upon deformation need to be skin passed or temper rolled to provide a source of dislocations for continuous yielding behavior. Steels displaying a yield point are undesirable for many forming operations because of the formation of Lüders bands that blemish the surface.
Hot dip galvanized products are produced on lines that process either preannealed (batch annealed) or full hard coils. Lines for processing full hard coils incorporate an in-line annealing capability so that annealing and hot dip galvanizing can be accomplished in a single pass through the line. This in-line annealing, like continuous annealing of uncoated steel, generally results in slightly higher strength and slightly lower ductility than batch annealing. Maximum strip temperatures are below the A1 temperature for commercial quality steel, but temperatures in excess of 845 °C (1550 °F) are required for DQSK grades. Galvanizing of preannealed steel results in properties similar to those of ungalvanized material. The atmosphere in a continuous galvanizing line, in addition to protecting the sheet from oxidation, must remove any oxides present on the strip to promote metallurgical bonding between the steel and the zinc or zinc alloy. Tin mill products are distinguished from their cold-rolled sheet mill counterparts chiefly by the fact that they are produced in lighter gages (0.13 to 0.38 mm, or 0.005 to 0.015 in.) and by the fact that some of them are coated with tin or chromium and chromium oxide for corrosion resistance. The sequence used for processing single-reduced tin mill products is similar to that for cold-rolled sheet, that is, pickling, cold reducing, annealing, and temper rolling of hot-rolled coils. Double-reduced products are cold rolled an additional 30 to 40% following annealing (this step replaces temper rolling). Whereas much of the tonnage produced in tin mills is batch annealed, a considerable amount is continuous annealed (facilities for continuous annealing currently are more prevalent in tin mills than in sheet mills). Because tin mill products traditionally have been produced at facilities separate from sheet mills and because applications for these products are different from those for cold-rolled sheet, tin mill products have been assigned separate designations for indicating the mechanical properties developed during annealing. A list of these temper designations is given in Table 5.
Open-coil annealing, which is done in batch furnaces, involves loose rewinding of a cold-reduced coil to provide open spaces between successive laps. This allows the controlled atmosphere gases to be drawn between the laps, providing faster and more uniform heating and cooling than are obtained with tightly wound coils. In addition, by control of the hydrogen content and dew point of the atmosphere, decarburizing conditions can be established. The carbon content of the steel can thereby be reduced to low levels for such materials as enameling steel and electrical steel.
Loose rewinding of coils for open-coil annealing is done on a turntable having a vertical mandrel. As the coil is wound, a twisted wire spacer is inserted between the laps. This spacer remains in the coil during annealing and is removed after the coil has been removed from the furnace. The coil is then tightly rewound and is ready for temper rolling.
Annealing of Steel Forgings
Annealing of forgings is most often performed to facilitate some subsequent operation, usually machining or cold forming. The type of annealing required is determined by the kind and amount of machining or cold forming to be done as well as the type of material involved. For some processes it is essential that the microstructure be spheroidal, whereas for others spheroidal structures may not be necessary or even desirable.
Annealing of Forgings for Machinability. In many cases, a structure suitable for machining can be developed in low-carbon steel forgings by transferring the forgings directly from the forging operation to a furnace heated to a proper transformation temperature, holding them at this temperature for a time sufficiently long to permit all the austenite to transform, then cooling in air. In this process, the effective austenitizing temperature is the finishing temperature of forging, not the initial forging temperature. This process is capable of producing reasonably uniform structures in forgings of uniform sections. However, in forgings shaped such that some portions are cooler than others, this difference in finishing temperature will cause the structures to be dissimilar. This process generally will not produce a spheroidal structure except in high-alloy steels containing large amounts of carbide-forming elements. If a lamellar structure is suitable for subsequent operations, however, this process can minimize the energy usage and lower costs by reducing processing and handling time.
In many instances where the product or subsequent process requires a more consistent hardness, forgings can be subcritical annealed by heating to a temperature between 10 and 20 °C (20 and 40 °F) below Ae1, holding sufficiently long (determined by degree of softening required), and then cooling in air (or equivalent). Care should be taken to maintain the temperature below Ae1 to prevent formation of austenite, which would require a much lower cooling rate. In forgings produced from higher-carbon steels with or without significant amounts of alloying elements, a spheroidal structure generally is preferable for high-speed machining operations. Direct transfer of high-carbon steel forgings to a furnace for transformation sometimes can be used as the preliminary step of an annealing cycle and as a means of preventing the possibility of cracking in deep-hardening steel parts, but seldom will produce satisfactory properties alone. Most annealing of high-carbon steel forgings is done either in a batch furnace or in a continuous tray pusher furnace. Typical schedules for spheroidizing 52100 steel in a batch furnace are as follows:
· Austenitize by holding at least 2 h at 790 °C (1450 °F), furnace cool at 17 °C/h (30 °F/h) to 595 °C (1100 °F), then air cool
· Austenitize by holding at least 2 h at 790 °C (1450 °F), cool as rapidly as practical to 750 °C (1380 °F), cool at 6 °C/h (10 °F/h) to 675 °C (1250 °F), then air cool
· Austenitize by holding at least 2 h at 790 °C (1450 °F), cool as rapidly as practical to 690 °C (1275 °F), transform isothermally by holding at this temperature for 16 h, then air cool
In all instances, the load should be distributed to promote uniform heating and cooling. Use of circulating fans in the furnace chamber will greatly aid in producing a product that is uniform in both hardness and microstructure. A typical continuous furnace for annealing steel forgings might consist of five or six zones. An example of a specific spheroidize annealing treatment in such a furnace is given in the next section.
Annealing of Forgings for Cold Forming and Re-Forming. If a steel forging or blank requires further cold forming, it may be necessary to soften it in order to enhance its plastic-flow characteristics. In general, this type of annealing is done only to the extent that the forming operation requires, that is, to satisfy dimensional, mechanical, and tool-life requirements, as well as to prevent cracking and splitting. Much intermediate annealing is done successfully, but cold forming processes are best performed on parts with totally spheroidized microstructures, especially for parts made of high-carbon steels. In one plant, both 5160 and 52100 steels have been successfully spheroidized with a common cycle in a six-zone tray pusher furnace. In this cycle, the temperatures in the six zones are 750, 750, 705, 695, 695, and 680 °C (1380, 1380, 1300, 1280, 1280, and 1260 °F). Time in each zone is 150 min. This process yields 5160 steel forgings with hardnesses of 170 to 190 RB and 52100 steel parts with hardnesses of 175 to 195 RB, both suitable for cold or warm restrike operations. In another cold-forming plant, 15B35 steel is processed in either a continuous roller hearth furnace or a bell furnace depending on the severity of the cold-forging operation. The continuous furnace is a two-zone furnace with zone temperatures at 750 °C (1380 °F) and 695 °C (1280 °F). Annealing time in each zone is between 90 to 120 min. The parts then enter a water-cooled cooling bed and exit the furnace at about 260 °C (500 °F). Only a partially spheroidized structure is obtainable in this furnace. If a nearly full-spheroidized structure is required, bell furnaces are used (see Fig. 13). A typical cycle for a 4500 kg (10,000 lb) load involves soaking at 760 °C (1400 °F) for 8 h followed by a slow cool to 675 °C (1250 °F) and a rapid cool afterwards.
A commercial heat treater uses a further variation of the processing cycle in its bell furnaces. The cycle for a 14,000 kg (31,000 lb) load involves soaking at 765 °C (1410 °F) for 24 h, furnace cooling to 675 °C (1250 °F), and holding at that temperature for 16 h followed by a rapid cooling.
Low-carbon steels generally can be cold formed successfully after being heated to temperatures near A1 and then being cooled through 675 °C (1250 °F) at a controlled rate. In one plant, 5120 steel annealed 1 to 2 h at 745 °C (1375 °F) and slow cooled has been cold formed successfully. Large quantities of 1008, 1513, 1524, 8620, and 8720 steels are being cold formed after annealing cycles consisting of 1 to 6 h at 720 °C (1325 °F) followed by slow cooling. The severity of the forming operation, as well as the grade of steel and history of the part, determines the extent of annealing required. Batch furnaces, continuous tray pusher furnaces, and continuous belt furnaces are being used successfully to perform these types of annealing operations on low-carbon steels.
Any part that contains significant stresses resulting from cold forming or restrike operations should be reviewed for some type of stress-relief process. Stress relieving usually is done by means of time-temperature cycles that result in slight reductions in hardness. These cycles often consist of 1 h at 425 to 675 °C (800 to 1250 °F).
Annealing to Obtain Pearlitic Microstructures. Forgings--especially plain and alloy high-carbon steel forgings--are isothermally annealed to produce a pearlitic microstructure that is preferred for a subsequent process. In steels that are to be induction hardened, for example, the carbide distribution of a fine pearlitic structure offers excellent preparation for optimum control in selective hardening while producing a reasonably machinable core structure. Isothermal annealing to obtain line pearlite can be performed in batch or continuous furnaces; however, temperature control and uniformity are more critical than in conventional slow cooling cycles because a particular microstructure and a particular hardness level usually are desired. In one plant, a continuous belt-type furnace is used for isothermal annealing of 1070 steel forgings. The forgings are uniformly heated for 30 min at 845 °C (1550 °F), cooled to 675 °C (1250 °F), and held for 20 min, then rapidly cooled. The microstructure produced is essentially fine lamellar pearlite with a hardness of 219 to 228 HB. The hardness and the structure can be modified by adjusting the transformation temperature.
Annealing of Bar, Rod, and Wire
Significant tonnages of bar, rod, and wire are subjected to thermal treatments that decrease hardness and prepare the material for subsequent cold working and/or machining. For low-carbon steels (up to 0.20% C), short-time subcritical annealing often is sufficient for preparing the material for further cold working. Steels with higher carbon and alloy contents require spheroidizing to impart maximum ductility.
Most carbon and alloy steel coiled products can be successfully spheroidized. In batch annealing, it is helpful to use higher-than-normal temperatures (for example, 650 °C, or 1200 °F) during initial heating for purging because the higher initial temperature promotes a lower temperature gradient in the charge during subsequent heating into the temperature range between A1 and A3. Use of a higher purge temperature also promotes agglomeration of the carbides in the steel, which makes them more resistant to dissolution in the austenite when the charge temperature is finally elevated. These undissolved carbides will be conducive to the formation of a spheroidal rather than a lamellar structure when transformation is complete.
A knowledge of the temperature distribution in the furnace and in the load can be a major factor in achieving a good, consistent response to spheroidization. Temperature distribution and control are much more critical in batch and vacuum furnaces, which may handle loads of up to 27 Mg (30 tons), than in continuous furnaces, in which loads of only 900 to 1800 kg (2000 to 4000 lb) may be transferred from zone to zone. Test thermocouples should be placed strategically at the top, middle, and bottom (inside and outside) of the charge during development of cycles. In spheroidizing, to minimize formation of pearlite on cooling, it is important to ensure that no part of the charge be allowed to approach A3. Conversely, if temperatures only slightly above A1 are used and temperature controls are inaccurate because of poor placement of thermocouples, it is probable that the A1 temperature will not be attained and that no austenitization will occur.
Table 6 gives typical mechanical properties that can be obtained in hypoeutectoid plain carbon steels by spheroidizing. Recommended temperatures and times for lamellar and spheroidize annealing of hypoeutectoid alloy steels are presented in Table 4.
Prior cold working increases the degree of spheroidization and provides even greater ductility. For example, 4037 steel in the as-rolled condition normally can be spheroidized to a tensile strength of about 515 MPa (75 ksi). If, however, the material is drawn 20% and then spheroidized (referred to as spheroidize annealed in-process), the resulting tensile strength will be around 470 MPa (68 ksi).
Although prior cold work can enhance response to annealing, caution must be observed in spheroidizing cold-worked plain carbon steels with 0.20% C or less. Unless a reduction of at least 20% is applied, severe grain coarsening may be observed after spheroidizing. Such grain coarsening is the result of a unique critical combination of strain and annealing temperature for the particular steel and may severely impair subsequent performance.
In the wire industry, a wide variety of in-process annealing operations have been evolved for rendering coiled material suitable for further processing that may require formability, drawability, machinability, or a combination of these characteristics. One large wire mill reports current use of 42 separate and distinct annealing cycles, the majority of which represent compromises between practical considerations and optimum properties. For example, annealing temperatures below those that might yield optimum softness sometimes must be used in order to preclude scaling of wire coils, which often can occur even in controlled-atmosphere furnaces. Even slight scaling may cause the coil wraps to stick together, which can impede coil payoff in subsequent operations.
Some of the terms used to describe various in-process annealing treatments are in common usage throughout the wire industry, whereas others have been developed within specific plants or mills. No attempt will be made here to list or define all the names that refer to specific treatments.
"Patenting" is a special form of annealing that is unique to the rod and wire industry. In this process, which usually is applied to medium- and higher-carbon grades of steel, rod or wire products are uncoiled, and the strands are delivered to an austenitizing station. The strands are then cooled rapidly from above A3 in a molten medium-- usually lead at about 540 °C (1000 °F)--for a period of time sufficient to allow complete transformation to a fine pearlitic structure. Both salt baths and fluidized beds have also been used for this purpose. This treatment increases substantially the amount of subsequent wiredrawing reduction that the product can withstand and permits production of high-strength wire. Successive drawing and patenting steps may be employed if necessary, in order to obtain the desired size and strength level.
Austenitizing for patenting can be accomplished in oil, gas, or electric furnaces; in high-temperature lead or salt baths; or by induction or direct resistance heating. As an alternative to quenching in molten lead, continuous air cooling often is employed. Such air patenting is less expensive than lead patenting but results in coarser pearlite and often more proeutectoid ferrite, a microstructure that is less desirable from the standpoint of drawing high-strength wire.
Plate products are occasionally annealed to facilitate forming or machining operations. Annealing of plate usually is done at subcritical temperatures, and long annealing times generally are avoided. Maintaining adequate flatness can be a significant problem in annealing of large plates.
Annealing of Tubular Products
Tubular products known as mechanical tubing are used in a variety of applications that can involve machining or forming. For these products, which are made from various grades of steel, annealing is a common treatment. In most annealing cycles, subcritical temperatures and short annealing times are used to reduce hardness to the desired level. High-carbon grades, such as 52100, generally are spheroidized to facilitate machining. Tubular products manufactured in pipe mills are rarely annealed. These products normally are used in the as-rolled, the normalized, or the quenched and tempered condition.
ANNEALING OF STEEL SHEET
products, a process often (but not solely) used to produce a recrystallized ferrite microstructure after cold rolling, is performed on a commercial scale by either batch annealing or continuous annealing. In batch annealing, multiple coils of sheet are placed under a cover with a reducing atmosphere (Fig. 1a) and heated for a time period that may involve days (see Fig. 1b). In contrast, continuous annealing of sheet involves the rapid passage of uncoiled sheet through heating and cooling equipment (Fig. 1a).
In addition to the obvious differences in equipment, batch and continuous annealing have important differences in heating and cooling profiles (Fig. 1b). The very large mass of steel heats and cools very slowly during batch annealing, and the process requires several days for completion. Annealed grain sizes are coarse, and the slow cooling rates ensure that all carbon dissolved during annealing precipitates upon cooling. Thus excellent ductility results, although some nonuniformity develops because the inside and outside parts of a coil experience different thermal histories. During continuous annealing, uncoiled steel sheet is passed though a two-stage furnace for times on the order of a few minutes. The first stage heats the steel and accomplishes recrystallization, while the second stage heats at a lower temperature to overage the steel and remove carbon from the solution effectively. Without this step, the thin sheet would cool too rapidly and retain carbon in solution. This carbon would eventually cause strain or quench aging and reduce sheet formability. There are several processing approaches to overaging, some of which are discussed in this article. For several decades, continuous-annealing lines have been widely used for the production of such sheet steel products as hot-dip galvanized steels, tinplate, nonoriented electrical steels, and stainless steels. The thermal profile on these lines generally involves short-time annealing followed by relatively slow cooling (~10 °C/s, or 20 °F/s) to ambient temperature with no inline overaging. Since the late 1970s, however, continuous-annealing technology, in conjunction with modern steelmaking and upstream processing facilities, has used rapid cooling and in-line overaging to enable the production of sheet steels for very demanding automotive and appliance applications.
The metallurgical advantages of continuous annealing over conventional batch annealing include improved product uniformity, surface cleanliness and shape, and the versatility to produce a wide range of steel grades.
Process Description
The modern continuous-annealing lines combine several processes. At the entry end, the uncoiled sheet is chemically and/or electrolytically cleaned and rapidly heated to an annealing temperature between 675 and 850 °C (1250 and 1550 °F). The cold-rolled sheet is "soaked" for an annealing time on the order of about 1 min and is then subjected to cooling and overaging (tempering) schedules such as those shown in Fig. 2. These stages in the heat treatment are discussed below. In most cases, a continuous-annealing line also includes a stage for tension leveling or temper rolling.
The heating and soaking/annealing stage provides recrystallization of the cold-rolled structure and achieves some degree of grain growth. The soaking temperature, which can range from 675 to 850 °C (1250 to 1560 °F), is generally above the A1 temperature. The low end of the annealing temperature range (675 °C, or 1250 °F) is used for commercial-quality (CQ) product
Friday, August 24, 2007
properties of c-40 steel
Component Wt. %
C 0.37 - 0.44
Fe 98.6 - 99
Mn 0.6 - 0.9
P Max 0.04
S Max 0.05
Material Notes:
Typical uses include machine, plow, and carriage bolts, tie wire, cylinder head studs, and machined parts, U-bolts, concrete reinforcing rods, forgings, and non-critical springs.
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Physical Properties Metric English Comments
Density 7.845 g/cc 0.283 lb/in³ Chemical composition of 0.435% C, 0.69% Mn, 0.20% Si, annealed at 860°C (1580°F).
Mechanical Properties
Hardness, Brinell 170 170
Hardness, Knoop 191 191 Converted from Brinell hardness.
Hardness, Rockwell B 86 86 Converted from Brinell hardness.
Hardness, Vickers 178 178 Converted from Brinell hardness.
Tensile Strength, Ultimate 620 MPa 89900 psi
Tensile Strength, Yield 550 MPa 79800 psi
Elongation at Break 12 % 12 % in 50 mm
Reduction of Area 35 % 35 %
Modulus of Elasticity 200 GPa 29000 ksi Typical for steel
Bulk Modulus 140 GPa 20300 ksi Typical for steels
Poisson's Ratio 0.29 0.29 Typical For Steel
Izod Impact 49 J 36.1 ft-lb as rolled, 45 J (33 ft-lb) annealed at 790°C (1450°F), 65 J (48 ft-lb) normalized at 900°C (1650°F)
Shear Modulus 80 GPa 11600 ksi Typical for steels
Electrical Properties
Electrical Resistivity 1.71e-005 ohm-cm 1.71e-005 ohm-cm 20°C (68°F)
Electrical Resistivity at Elevated Temperature 0.0001111 ohm-cm 0.0001111 ohm-cm 800°C (1470°F)
Electrical Resistivity at Elevated Temperature 0.0001149 ohm-cm 0.0001149 ohm-cm 900°C (1650°F)
Electrical Resistivity at Elevated Temperature 0.0001179 ohm-cm 0.0001179 ohm-cm 1000°C
Electrical Resistivity at Elevated Temperature 2.21e-005 ohm-cm 2.21e-005 ohm-cm 100°C (212°F)
Electrical Resistivity at Elevated Temperature 2.96e-005 ohm-cm 2.96e-005 ohm-cm 200°C (390°F)
Electrical Resistivity at Elevated Temperature 4.93e-005 ohm-cm 4.93e-005 ohm-cm 400°C (750°F)
Electrical Resistivity at Elevated Temperature 7.63e-005 ohm-cm 7.63e-005 ohm-cm 600°C (1110°F)
Electrical Resistivity at Elevated Temperature 9.32e-005 ohm-cm 9.32e-005 ohm-cm 700°C (1290°F)
Thermal Properties
CTE, linear 20°C 11.3 µm/m-°C 6.28 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu.; 20-100°C (68-212°F)
CTE, linear 20°C 12.1 µm/m-°C 6.72 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu.; 20-200°C (68-390°F)
CTE, linear 250°C 12.2 µm/m-°C 6.78 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-300°C (68-570°F)
CTE, linear 250°C 13.3 µm/m-°C 7.39 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-400°C (68-750°F)
CTE, linear 500°C 13.9 µm/m-°C 7.72 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-500°C (68-930°F)
CTE, linear 500°C 14.2 µm/m-°C 7.89 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-600°C (68-1110°F)
CTE, linear 500°C 14.8 µm/m-°C 8.22 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-700°C (68-1290°F)
CTE, linear 1000°C 14.7 µm/m-°C 8.17 µin/in-°F Typical steel
Specific Heat Capacity 0.486 J/g-°C 0.116 BTU/lb-°F 50-100°C (122-212°F)
Specific Heat Capacity at Elevated Temperature 0.515 J/g-°C 0.123 BTU/lb-°F 150-200°C (300-390°F)
Specific Heat Capacity at Elevated Temperature 0.528 J/g-°C 0.126 BTU/lb-°F 200-250°C (390-480°F)
Specific Heat Capacity at Elevated Temperature 0.548 J/g-°C 0.131 BTU/lb-°F 250-300°C (480-570°F)
Specific Heat Capacity at Elevated Temperature 0.569 J/g-°C 0.136 BTU/lb-°F 300-350°C (570-660°F)
Specific Heat Capacity at Elevated Temperature 0.586 J/g-°C 0.14 BTU/lb-°F 350-400°C (660-750°F)
Specific Heat Capacity at Elevated Temperature 0.624 J/g-°C 0.149 BTU/lb-°F 750-800°C (1380-1470°F)
Specific Heat Capacity at Elevated Temperature 0.649 J/g-°C 0.155 BTU/lb-°F 450-500°C (750-930°F)
Specific Heat Capacity at Elevated Temperature 0.708 J/g-°C 0.169 BTU/lb-°F 550-600°C (1020-1110°F)
Specific Heat Capacity at Elevated Temperature 0.77 J/g-°C 0.184 BTU/lb-°F 650-700°C (1200-1290°F)
Specific Heat Capacity at Elevated Temperature 1.583 J/g-°C 0.378 BTU/lb-°F 700-750°C (1290-1380°F)
Thermal Conductivity 50.7 W/m-K 352 BTU-in/hr-ft²-°F 100°C (212°F)
Thermal Conductivity 51.9 W/m-K 360 BTU-in/hr-ft²-°F 0°C
Thermal Conductivity at Elevated Temperature 24.7 W/m-K 171 BTU-in/hr-ft²-°F 800°C
Thermal Conductivity at Elevated Temperature 29.8 W/m-K 207 BTU-in/hr-ft²-°F 1200°C (2190°F)
Thermal Conductivity at Elevated Temperature 30.1 W/m-K 209 BTU-in/hr-ft²-°F 700°C (1290°F)
Thermal Conductivity at Elevated Temperature 32.9 W/m-K 228 BTU-in/hr-ft²-°F 1000°C (1830°F)
Thermal Conductivity at Elevated Temperature 33.9 W/m-K 235 BTU-in/hr-ft²-°F 600°C (1110°F)
Thermal Conductivity at Elevated Temperature 38.2 W/m-K 265 BTU-in/hr-ft²-°F 500°C (930°F)
Thermal Conductivity at Elevated Temperature 41.7 W/m-K 289 BTU-in/hr-ft²-°F 400°C (750°F)
Thermal Conductivity at Elevated Temperature 45.7 W/m-K 317 BTU-in/hr-ft²-°F 300°C (570°F)
Thermal Conductivity at Elevated Temperature 48.1 W/m-K 334 BTU-in/hr-ft²-°F 200°C (390°F)
C 0.37 - 0.44
Fe 98.6 - 99
Mn 0.6 - 0.9
P Max 0.04
S Max 0.05
Material Notes:
Typical uses include machine, plow, and carriage bolts, tie wire, cylinder head studs, and machined parts, U-bolts, concrete reinforcing rods, forgings, and non-critical springs.
Click here to view available vendors for this material.
Physical Properties Metric English Comments
Density 7.845 g/cc 0.283 lb/in³ Chemical composition of 0.435% C, 0.69% Mn, 0.20% Si, annealed at 860°C (1580°F).
Mechanical Properties
Hardness, Brinell 170 170
Hardness, Knoop 191 191 Converted from Brinell hardness.
Hardness, Rockwell B 86 86 Converted from Brinell hardness.
Hardness, Vickers 178 178 Converted from Brinell hardness.
Tensile Strength, Ultimate 620 MPa 89900 psi
Tensile Strength, Yield 550 MPa 79800 psi
Elongation at Break 12 % 12 % in 50 mm
Reduction of Area 35 % 35 %
Modulus of Elasticity 200 GPa 29000 ksi Typical for steel
Bulk Modulus 140 GPa 20300 ksi Typical for steels
Poisson's Ratio 0.29 0.29 Typical For Steel
Izod Impact 49 J 36.1 ft-lb as rolled, 45 J (33 ft-lb) annealed at 790°C (1450°F), 65 J (48 ft-lb) normalized at 900°C (1650°F)
Shear Modulus 80 GPa 11600 ksi Typical for steels
Electrical Properties
Electrical Resistivity 1.71e-005 ohm-cm 1.71e-005 ohm-cm 20°C (68°F)
Electrical Resistivity at Elevated Temperature 0.0001111 ohm-cm 0.0001111 ohm-cm 800°C (1470°F)
Electrical Resistivity at Elevated Temperature 0.0001149 ohm-cm 0.0001149 ohm-cm 900°C (1650°F)
Electrical Resistivity at Elevated Temperature 0.0001179 ohm-cm 0.0001179 ohm-cm 1000°C
Electrical Resistivity at Elevated Temperature 2.21e-005 ohm-cm 2.21e-005 ohm-cm 100°C (212°F)
Electrical Resistivity at Elevated Temperature 2.96e-005 ohm-cm 2.96e-005 ohm-cm 200°C (390°F)
Electrical Resistivity at Elevated Temperature 4.93e-005 ohm-cm 4.93e-005 ohm-cm 400°C (750°F)
Electrical Resistivity at Elevated Temperature 7.63e-005 ohm-cm 7.63e-005 ohm-cm 600°C (1110°F)
Electrical Resistivity at Elevated Temperature 9.32e-005 ohm-cm 9.32e-005 ohm-cm 700°C (1290°F)
Thermal Properties
CTE, linear 20°C 11.3 µm/m-°C 6.28 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu.; 20-100°C (68-212°F)
CTE, linear 20°C 12.1 µm/m-°C 6.72 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu.; 20-200°C (68-390°F)
CTE, linear 250°C 12.2 µm/m-°C 6.78 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-300°C (68-570°F)
CTE, linear 250°C 13.3 µm/m-°C 7.39 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-400°C (68-750°F)
CTE, linear 500°C 13.9 µm/m-°C 7.72 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-500°C (68-930°F)
CTE, linear 500°C 14.2 µm/m-°C 7.89 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-600°C (68-1110°F)
CTE, linear 500°C 14.8 µm/m-°C 8.22 µin/in-°F Composition of 0.40% C, 0.11% Mn, 0.01% P, 0.03% S, 0.03% Si, 0.03% Cu; 20-700°C (68-1290°F)
CTE, linear 1000°C 14.7 µm/m-°C 8.17 µin/in-°F Typical steel
Specific Heat Capacity 0.486 J/g-°C 0.116 BTU/lb-°F 50-100°C (122-212°F)
Specific Heat Capacity at Elevated Temperature 0.515 J/g-°C 0.123 BTU/lb-°F 150-200°C (300-390°F)
Specific Heat Capacity at Elevated Temperature 0.528 J/g-°C 0.126 BTU/lb-°F 200-250°C (390-480°F)
Specific Heat Capacity at Elevated Temperature 0.548 J/g-°C 0.131 BTU/lb-°F 250-300°C (480-570°F)
Specific Heat Capacity at Elevated Temperature 0.569 J/g-°C 0.136 BTU/lb-°F 300-350°C (570-660°F)
Specific Heat Capacity at Elevated Temperature 0.586 J/g-°C 0.14 BTU/lb-°F 350-400°C (660-750°F)
Specific Heat Capacity at Elevated Temperature 0.624 J/g-°C 0.149 BTU/lb-°F 750-800°C (1380-1470°F)
Specific Heat Capacity at Elevated Temperature 0.649 J/g-°C 0.155 BTU/lb-°F 450-500°C (750-930°F)
Specific Heat Capacity at Elevated Temperature 0.708 J/g-°C 0.169 BTU/lb-°F 550-600°C (1020-1110°F)
Specific Heat Capacity at Elevated Temperature 0.77 J/g-°C 0.184 BTU/lb-°F 650-700°C (1200-1290°F)
Specific Heat Capacity at Elevated Temperature 1.583 J/g-°C 0.378 BTU/lb-°F 700-750°C (1290-1380°F)
Thermal Conductivity 50.7 W/m-K 352 BTU-in/hr-ft²-°F 100°C (212°F)
Thermal Conductivity 51.9 W/m-K 360 BTU-in/hr-ft²-°F 0°C
Thermal Conductivity at Elevated Temperature 24.7 W/m-K 171 BTU-in/hr-ft²-°F 800°C
Thermal Conductivity at Elevated Temperature 29.8 W/m-K 207 BTU-in/hr-ft²-°F 1200°C (2190°F)
Thermal Conductivity at Elevated Temperature 30.1 W/m-K 209 BTU-in/hr-ft²-°F 700°C (1290°F)
Thermal Conductivity at Elevated Temperature 32.9 W/m-K 228 BTU-in/hr-ft²-°F 1000°C (1830°F)
Thermal Conductivity at Elevated Temperature 33.9 W/m-K 235 BTU-in/hr-ft²-°F 600°C (1110°F)
Thermal Conductivity at Elevated Temperature 38.2 W/m-K 265 BTU-in/hr-ft²-°F 500°C (930°F)
Thermal Conductivity at Elevated Temperature 41.7 W/m-K 289 BTU-in/hr-ft²-°F 400°C (750°F)
Thermal Conductivity at Elevated Temperature 45.7 W/m-K 317 BTU-in/hr-ft²-°F 300°C (570°F)
Thermal Conductivity at Elevated Temperature 48.1 W/m-K 334 BTU-in/hr-ft²-°F 200°C (390°F)
Thursday, August 9, 2007
Annealing
Annealing
Abstract:
The purpose of annealing may involve one or more of the following aims:
1. To soften the steel and to improve machinability.
2. To relieve internal stresses induced by some previous treatment (rolling, forging,uneven cooling).
3. To remove coarseness of grain.
The treatment is applied to forgings, cold-worked sheets and wire, and castings. The operation consists of:
a. heating the steel to a certain temperature,
b. "soaking" at this temperature for a time sufficient to allow the necessary changes to occur,
c. cooling at a predetermined rate.
The purpose of annealing may involve one or more of the following aims:
To soften the steel and to improve machinability.
To relieve internal stresses induced by some previous treatment (rolling, forging, uneven cooling).
To remove coarseness of grain.
The treatment is applied to forgings, cold-worked sheets and wire, and castings. The operation consists of:
heating the steel to a certain temperature,
"soaking" at this temperature for a time sufficient to allow the necessary changes to occur,
cooling at a predetermined rate.
Sub-critical Anneal
It is not always necessary to heat the steel into the critical range. Mild steel products which have to be repeatedly cold worked in the processes of manufacture are softened by annealing at 500° to 650°C for several hours. This is known as "process" or "close" annealing, and is commonly employed for wire and sheets. The recrystallisation temperature of pure iron is in the region of 500°C consequently the higher temperature of 650°C brings about rapid recrystallisation of the distorted ferrite Since mild steel contains only a small volume of strained pearlite a high degree of softening is induced. As shown, Fig. 1b illustrates the structure formed consisting of the polyhedral ferrite with elongated pearlite (see also Fig. 2).
Prolonged annealing induces greater ductility at the expense of strength, owing to the tendency of the cementite in the strained pearlite to "ball-up" or spheroidise, as illustrated in Fig. 1c. This is known as "divorced pearlite". The ferrite grains also become larger, particularly if the metal has been cold worked a critical amount. A serious embrittlement sometimes arises after prolonged treatment owing to the formation of cementitic films at the ferrite boundaries. With severe forming operations, cracks are liable to start at these cementite membranes.
The modern tendency is to use batch or continuous annealing furnaces with an inert purging gas. Batch annealing usually consists of 24-30 hrs 670°C, soak 12 hrs, slow cool 4-5 days. Open coil annealing consists in recoiling loosely with controlled space between wraps and it reduces stickers and discoloration. Continuous annealing is used for thin strip (85% Red) running at about 400 m/min. The cycle is approximately up to 660°C 20 sec, soak and cool 30-40 sec. There is little chance for grain growth and it produces harder and stiffer strip; useful for cans and panelling.
"Double reduced" steel is formed by heavy reduction (~50%) after annealing but it suffers from directionality. This can be eliminated by heating between 700-920°C and rapidly quenching.
Full Anneal and Normalising Treatments
For steels with less than 0,9% carbon both treatments consist in heating to about 25-50°C above the upper critical point indicated by the Fe-Fe3C equilibrium diagram (Fig. 3). For higher carbon steels the temperature is 50°C above the lower critical point.
Average annealing and hardening temperatures are:
Carbon, % 0.1 0.2 0.3 0.5 0.7 0.9 to 1.3
Avg.temp. °C 910 860 830 810 770 760
These temperatures allow for the effects of slight variations in the impurities present and also the thermal lag associated with the critical changes. After soaking at the temperature for a time dependent on the thickness of the article, the steel is very slowly cooled. This treatment is known as full annealing, and is used for removing strains from forgings and castings, improving machinability and also when softening and refinement of structure are both required.
Normalising differs from the full annealing in that the metal is allowed to cool in still air. The structure and properties produced, however, varying with the thickness of metal treated. The tensile strength, yield point, reduction of area and impact value are higher than the figures obtained by annealing.
Changes on Annealing
Consider the heating of a 0,3% carbon steel. At the lower critical point (Ac1) each "grain" of pearlite changes to several minute austenite crystals and as the temperature is raised the excess ferrite is dissolved, finally disappearing at the upper critical point (Ac3), still with the production of fine austenite crystals. Time is necessary for the carbon to become uniformly distributed in this austenite. The properties obtained subsequently depend on the coarseness of the pearlite and ferrite and their relative distribution. These depend on:
a) the size of the austenite grains; the smaller their size the better the distribution of the ferrite and pearlite.
b) the rate of cooling through the critical range, which affects both the ferrite and the pearlite.
As the temperature is raised above Ac3 the crystals increase in size. On a certain temperature the growth, which is rapid at first, diminishes. Treatment just above the upper critical point should be aimed at, since the austenite crystals are then small.
By cooling slowly through the critical range, ferrite commences to deposit on a few nuclei at the austenite boundaries. Large rounded ferrite crystals are formed, evenly distributed among the relatively coarse pearlite. With a higher rate of cooling, many ferrite crystals are formed at the austenite boundaries and a network structure of small ferrite crystals is produced with fine pearlite in the centre.
Overheated, Burnt and Underannealed Structures
When the steel is heated well above the upper critical temperature large austenite crystals form. Slow cooling gives rise to the Widmanstätten type of structure, with its characteristic lack of both ductility and resistance to shock. This is known as an overheated structure, and it can be refined by reheating the steel to just above the upper critical point. Surface decarburisation usually occurs during the overheating.
During the Second World War, aircraft engine makers were troubled with overheating (above 1250°C) in drop-stampings made from alloy steels. In the hardened and tempered condition the fractured surface shows dull facets. The minimum overheating temperature depends on the "purity" of the steel and is substantially lower in general for electric steel than for open-hearth steel. The overheated structure in these alloy steels occurs when they are cooled at an intermediate rate from the high temperature. At faster or slower rates the overheated structure may be eliminated. This, together with the fact that the overheating temperature is significantly raised in the presence of high contents of MnS and inclusions, suggests that this overheating is conected in some way with a diffusion and precipitation process, involving MnS. This type of overheating can occur in an atmosphere free from oxygen, thus emphasising the difference between overheating and burning.
As the steel approaches the solidus temperature, incipient fusion and oxidation take place at the grain boundaries. Such a steel is said to be burnt and it is characterised by the presence of brittle iron oxide films, which render the steel unfit for service, except as scrap for remelting.
Abstract:
The purpose of annealing may involve one or more of the following aims:
1. To soften the steel and to improve machinability.
2. To relieve internal stresses induced by some previous treatment (rolling, forging,uneven cooling).
3. To remove coarseness of grain.
The treatment is applied to forgings, cold-worked sheets and wire, and castings. The operation consists of:
a. heating the steel to a certain temperature,
b. "soaking" at this temperature for a time sufficient to allow the necessary changes to occur,
c. cooling at a predetermined rate.
The purpose of annealing may involve one or more of the following aims:
To soften the steel and to improve machinability.
To relieve internal stresses induced by some previous treatment (rolling, forging, uneven cooling).
To remove coarseness of grain.
The treatment is applied to forgings, cold-worked sheets and wire, and castings. The operation consists of:
heating the steel to a certain temperature,
"soaking" at this temperature for a time sufficient to allow the necessary changes to occur,
cooling at a predetermined rate.
Sub-critical Anneal
It is not always necessary to heat the steel into the critical range. Mild steel products which have to be repeatedly cold worked in the processes of manufacture are softened by annealing at 500° to 650°C for several hours. This is known as "process" or "close" annealing, and is commonly employed for wire and sheets. The recrystallisation temperature of pure iron is in the region of 500°C consequently the higher temperature of 650°C brings about rapid recrystallisation of the distorted ferrite Since mild steel contains only a small volume of strained pearlite a high degree of softening is induced. As shown, Fig. 1b illustrates the structure formed consisting of the polyhedral ferrite with elongated pearlite (see also Fig. 2).
Prolonged annealing induces greater ductility at the expense of strength, owing to the tendency of the cementite in the strained pearlite to "ball-up" or spheroidise, as illustrated in Fig. 1c. This is known as "divorced pearlite". The ferrite grains also become larger, particularly if the metal has been cold worked a critical amount. A serious embrittlement sometimes arises after prolonged treatment owing to the formation of cementitic films at the ferrite boundaries. With severe forming operations, cracks are liable to start at these cementite membranes.
The modern tendency is to use batch or continuous annealing furnaces with an inert purging gas. Batch annealing usually consists of 24-30 hrs 670°C, soak 12 hrs, slow cool 4-5 days. Open coil annealing consists in recoiling loosely with controlled space between wraps and it reduces stickers and discoloration. Continuous annealing is used for thin strip (85% Red) running at about 400 m/min. The cycle is approximately up to 660°C 20 sec, soak and cool 30-40 sec. There is little chance for grain growth and it produces harder and stiffer strip; useful for cans and panelling.
"Double reduced" steel is formed by heavy reduction (~50%) after annealing but it suffers from directionality. This can be eliminated by heating between 700-920°C and rapidly quenching.
Full Anneal and Normalising Treatments
For steels with less than 0,9% carbon both treatments consist in heating to about 25-50°C above the upper critical point indicated by the Fe-Fe3C equilibrium diagram (Fig. 3). For higher carbon steels the temperature is 50°C above the lower critical point.
Average annealing and hardening temperatures are:
Carbon, % 0.1 0.2 0.3 0.5 0.7 0.9 to 1.3
Avg.temp. °C 910 860 830 810 770 760
These temperatures allow for the effects of slight variations in the impurities present and also the thermal lag associated with the critical changes. After soaking at the temperature for a time dependent on the thickness of the article, the steel is very slowly cooled. This treatment is known as full annealing, and is used for removing strains from forgings and castings, improving machinability and also when softening and refinement of structure are both required.
Normalising differs from the full annealing in that the metal is allowed to cool in still air. The structure and properties produced, however, varying with the thickness of metal treated. The tensile strength, yield point, reduction of area and impact value are higher than the figures obtained by annealing.
Changes on Annealing
Consider the heating of a 0,3% carbon steel. At the lower critical point (Ac1) each "grain" of pearlite changes to several minute austenite crystals and as the temperature is raised the excess ferrite is dissolved, finally disappearing at the upper critical point (Ac3), still with the production of fine austenite crystals. Time is necessary for the carbon to become uniformly distributed in this austenite. The properties obtained subsequently depend on the coarseness of the pearlite and ferrite and their relative distribution. These depend on:
a) the size of the austenite grains; the smaller their size the better the distribution of the ferrite and pearlite.
b) the rate of cooling through the critical range, which affects both the ferrite and the pearlite.
As the temperature is raised above Ac3 the crystals increase in size. On a certain temperature the growth, which is rapid at first, diminishes. Treatment just above the upper critical point should be aimed at, since the austenite crystals are then small.
By cooling slowly through the critical range, ferrite commences to deposit on a few nuclei at the austenite boundaries. Large rounded ferrite crystals are formed, evenly distributed among the relatively coarse pearlite. With a higher rate of cooling, many ferrite crystals are formed at the austenite boundaries and a network structure of small ferrite crystals is produced with fine pearlite in the centre.
Overheated, Burnt and Underannealed Structures
When the steel is heated well above the upper critical temperature large austenite crystals form. Slow cooling gives rise to the Widmanstätten type of structure, with its characteristic lack of both ductility and resistance to shock. This is known as an overheated structure, and it can be refined by reheating the steel to just above the upper critical point. Surface decarburisation usually occurs during the overheating.
During the Second World War, aircraft engine makers were troubled with overheating (above 1250°C) in drop-stampings made from alloy steels. In the hardened and tempered condition the fractured surface shows dull facets. The minimum overheating temperature depends on the "purity" of the steel and is substantially lower in general for electric steel than for open-hearth steel. The overheated structure in these alloy steels occurs when they are cooled at an intermediate rate from the high temperature. At faster or slower rates the overheated structure may be eliminated. This, together with the fact that the overheating temperature is significantly raised in the presence of high contents of MnS and inclusions, suggests that this overheating is conected in some way with a diffusion and precipitation process, involving MnS. This type of overheating can occur in an atmosphere free from oxygen, thus emphasising the difference between overheating and burning.
As the steel approaches the solidus temperature, incipient fusion and oxidation take place at the grain boundaries. Such a steel is said to be burnt and it is characterised by the presence of brittle iron oxide films, which render the steel unfit for service, except as scrap for remelting.
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