DUCTILE IRON DATA FOR DESIGN ENGINEERS
SECTION III. ENGINEERING DATA (part 2)
Figures 3.26, 3.29 and 3.30 indicate that matrix type and related mechanical properties, especially tensile strength and hardness, exert considerable influence on fatigue strength. However, the decrease in endurance ratio with increasing tensile strength in Figures 3.26 indicates that increasing the tensile strength of Ductile Iron does not provide a proportionate increase in fatigue strength. Figure 3.29 shows that, for constant nodule size, fatigue strength increases with Vickers micro-hardness number, reaching a maximum at a hardness value of 500. Examination of Figure 3.30 reveals a significant influence of matrix micro-hardness on fatigue strength at low inclusion levels, which declines as the volume fraction of inclusions increases.
Figure 3.32, from Janowak, Alagarsamy and Venugopalan, indicates that there is a good correlation between fatigue strength and the calculated composite matrix micro hardness (CMMH). (See Figure 3.9 for a similar relationship between tensile properties and CMMH.) Figure 3.32 also includes the data of Sofue et al, from whose work Figures 3.29 and 3.30 are taken. The region marked "gear" in this Figure refers to data reported by Sofue et al on the successful performance of pearlitic Ductile Iron gears with induction hardened teeth. It is interesting to note that the fatigue performance of commercial Ductile Irons shown in Figure 3.32 is superior, at equal hardness, to that of the irons produced in the laboratory by Sofue et al (D1-D8). Janowak et al attributed the inferior performance of the laboratory irons to low alloy and residual element contents, and the quench and temper heat treatments used by Sofue et al to produce different matrix hardness levels. Nevertheless, Figure 3.32 confirms that a good correlation exists between matrix microhardness and fatigue strength and that the fatigue performance of Ductile Iron can be predicted using the calculated CMMH.
Because fatigue failures generally occur after a significant period of time has elapsed, fatigue behaviour can be degraded significantly by environments which accelerate crack initiation and growth. Figure 3.33 illustrates the reduction in fatigue strength resulting from exposure to water spray environments consisting or water and aqueous solutions of borax, sodium carbonate, and a soluble oil. In the most aggressive environment, borax, fatigue strength was reduced by 28 per cent. In accord with the time-dependent nature of corrosion-assisted fatigue, the effect of the corrosive environments decreased with decreasing fatigue life. Only potassium chromate, an inhibitor, prevented any significant loss in fatigue strength due to exposure to an aqueous environment. Chromate solutions are now considered to be toxic, and a combination of 0.5% sodium nitrate and 1% sodium silicate has been shown to be equally effective. Table 3.2 shows that spray coatings of zinc and aluminium provide excellent protection against corrosion fatigue of Ductile Iron by water and brine spray environments. Uncoated samples showed fatigue strength reductions of 1. 2 and 5.8 times respectively in water and brine sprays, while zinc- and aluminium-coated samples showed no loss of fatigue strength.
Table 3.2 Effect of environment and coatings on corrosion fatigue strength of pearlitic Ductile Iron.
In bending and torsional fatigue, in which cyclic stresses attain maximum values at the component surface, fatigue behaviour is strongly dependent upon surface geometry, residual stress conditions and material properties in the surface layer of the component. The use of adequate fillet radii, shot peening, surface rolling, flame and induction hardening and nitriding can significantly increase the fatigue limit of Ductile Iron components. These treatments, which will be discussed in more detail in Section IX, enhance fatigue resistance by 20 to 100 per cent by increasing the tensile strength and inducing compressive stresses in the surface layer of the component. In addition to improving surface stress conditions, shot peening also reduces the stress concentration effects of surface roughness.
Figure 3.34 illustrates the effect of different levels of shot peening intensity on the fatigue strength of pearlitic Ductile Iron with as-cast surfaces. Shot peening at the highest intensity level developed fatigue properties of the as-cast surfaces to within 6 per cent of those with defect-free machined surfaces.
Figure 3.35 illustrates the influence of surface rolling on the bending fatigue properties of ferritic and pearlitic grades of Ductile Iron. This Figure shows that v-notched samples, strengthened by rolling with a roller contoured to the notch geometry, had fatigue strengths from 58 to 73 per cent higher than the unnotched samples of the pearlitic and ferritic grades respectively. Table 3.3, which compares the reversed bending fatigue properties of different Ductile Iron crankshafts, confirms the significant strengthening effect of fillet rolling. Fillet rolling of ascast crankshafts increased fatigue strength from 30 ksi (207 MPa) to 83-97 ksi (572-669 MPa), an increase of 175-225 per cent over the as-cast pearlitic iron. This Table also documents the even greater benefits accruing from austempering and fillet rolling (see Section IV for more information on the fatigue properties of austempered Ductile Iron).
Surface hardening by flame or induction heating is used to improve the resistance of Ductile Iron to both normal and pitting fatigue failures. Conventional fatigue strength is improved by a combination of high surface hardness and compressive surface stresses, while pitting fatigue is reduced by the increased surface hardness. Molten salt cyaniding produces a two-layer "case" on Ductile Iron components which can result in increases in fatigue strengths from 63 to 80 per cent, as shown in Figure 3.36.
The design stress for fatigue should not exceed one-third of the fatigue limit measured under conditions that suitably replicate the stress environment of the application. That is, notched data should be used when unavoidable stress concentrations are present in the component, and bending, torsional and push-pull fatigue data should be used according to the type of cyclic stress encountered by the component. The fatigue strength of Ductile Iron is frequency sensitive, and test frequencies should not exceed those encountered when the component is in service. The fatigue strength of Ductile Iron, like many other cast materials, is also influenced by both the cast section size and the specimen size. Both of these factors should be considered when extrapolating laboratory fatigue data to actual components, although the one-third safety factor may be sufficient to compensate for any degradation in fatigue strength due to size factors. The fatigue strength of Ductile Iron can be optimized through a combination of production and design practices which result in the following component characteristics.
Thermal fatigue is a special type of fatigue in which thermal cycling produces stress/ strain cycles in the component through differential expansion and contraction resulting from temperature gradients. The severity of thermal fatigue increases with increased temperature, increased range over which the temperature is cycled and increased rates of heating and cooling. Material properties which contribute to good thermal fatigue resistance are: high thermal conductivity, low modulus of elasticity and high strength and ductility. For severe thermal fatigue conditions, the high thermal conductivity and low modulus of high carbon Gray Iron make this material superior to both conventional and alloyed ferritic Ductile Irons and Compacted Graphite (CG) Iron.
For medium severity thermal fatigue, ferritic Ductile Iron and CG Iron provide superior cracking resistance but may fail by distortion. Pearlitic and alloy Ductile Irons provide the best performance for low severity thermal fatigue conditions. Figure 3.37 shows the increasing superiority of ferritic, pearlitic and alloy Ductile Irons in the Buderus Test in which thermal fatigue resistance is ranked by measuring the number of cycles between 650oC (1200oF) and room temperature required to produce bridge cracking between two holes in the test specimen. Performance of exhaust manifolds follows closely the ranking shown in this Figure. Ferritic Ductile Iron exhaust manifolds have been used widely due to a combination of good thermal fatigue strength and resistance to graphitization. Recent demands for increased service temperatures have resulted in the use of "Si-Mo" Ductile Irons containing 4-5% Si and up to 1% Mo. The increased strength and oxidation resistance of these alloys have resulted in excellent performance at service temperatures up to 750oC (1380oF).
Ductile Iron, like most ferrous materials, exhibits fracture behaviour which varies according to composition, microstructure, temperature, strain rate, and stress state. At low temperatures, brittle failure occurs by the formation of cleavage cracks, producing a facetted, shiny fracture surface. Very little deformation is associated with this type of fracture, resulting in low absorption of energy and low toughness. As the temperature increases, producing a decrease in flow stress, failure occurs by plastic deformation, primarily by the formation, growth and coalescence of voids. The resultant fracture surface will be dull gray, and the energy absorbed will be high, meaning very good fracture toughness. Fracture in ferrous materials traditionally has been characterized according to appearance and absorbed energy, with a Nil-Ductility-Transition (NDT) temperature quoted to indicate the change from brittle to ductile behaviour. In addition to transition temperature, upper shelf energies were quoted to define toughness in the ductile fracture region.
The Charpy test has been used for many years to characterize both the transition temperature and fracture energy for Ductile Iron, and a large body of Charpy impact energy data has been accumulated. The Charpy test is a dynamic fracture test in which a notched (see Figure 3.38) or unnotched test piece is struck an impact blow by a swinging pendulum. The effect of the notch on the fracture behaviour of ferritic Ductile Iron is shown in Figure 3.39. The shape of the notch is also important and must be considered, "V" shaped notches being more severe and producing lower strengths than "U" notches. The complex, triaxial stress state and increased strain rate at the root of the notch combine to restrict plastic deformation, increasing the transition temperature by 110oF (60oC) and reducing the upper shelf energy by 75 per cent. The effect of strain rate on fracture behaviour is illustrated in Figure 3.40, in which the results of dynamic (impact) tests of pre-cracked, notched Charpy bars are compared to quasi-static (slow bend) test results. The increased loading rate of the impact test produced a 115oF (64oC) increase in the transition temperature. Figure 3.39 and Figure 3.40 highlight the sensitivity of fracture behaviour to test conditions and emphasize the strain rate sensitivity of Ductile Iron.
Recent advances in fracture mechanics have resulted in the use of the Dynamic Tear Test, ASTM E604, and the fracture toughness tests, ASTM E399 and ASTM E813,(see Figure 3.38 for sample geometry) to determine crack propagation properties, which are considered more relevant to the assessment of the flaw tolerance of a stressed component. This section will use data obtained from standard and modified Charpy tests and from dynamic tear and fracture toughness tests to characterize the fracture behaviour of Ductile Iron. The large body of standard Charpy data will be used to illustrate the relative effects of microstructure, composition, heat treatment and stress environment on fracture behaviour. Data from the other tests are offered to provide the designer with the quantitative information required to make materials selection and component design decisions. Again, fracture toughness informatio is more relevant and Ductile Iron compares well with steel in toughness levels where it is not shown to be as good with Charpy data.
The impact properties of Ductile Iron are influenced significantly by matrix microstructure. As shown in Figure 3.41, Ductile Irons with annealed ferritic matrices exhibit the lowest ductile-to-brittle transition temperature and highest upper shelf energy. Special, annealed or sub-critically annealed ferritic Ductile Irons with tensile strengths at or below 60,000 psi (414 MPa) are normally specified when very high notch ductility and good low temperature toughness are required. These irons have v-notched Charpy impact transition temperatures in the range 0oC to - 60o C (32oF to - 76o F), depending on heat treatment, composition and graphite properties (see Figures 3.42 , 3.43, 3.44, 3.45, and 3.46). These materials normally exhibit upper shelf energies in the range 16-24 joules (12-18 ft lbf) with room temperature values in excess of 16 Joules (12 ft lbf).
As-cast ferritic grades, and those with increasing percentages of pearlite, have increasingly higher transition temperatures and lower upper shelf energies. Generally, pearlitic grades of Ductile Iron are used because of their higher strengths in applications requiring only limited ductility and toughness and are generally not recommended for use in low temperature applications requiring impact resistance. However, in spite of apparently poor low temperature toughness, hundreds of ASTM Grade 100-70-03 gears have performed without problems in oilfield pumps operating at subzero temperatures in northern climates. Quenched and tempered martensitic Ductile Irons generally exhibit a combination of strength and low temperature toughness (see Figure 3.44) that is superior to those of pearlitic grades.
In addition to influencing microstructural characteristics such as ferrite: pearlite ratio and carbide content, composition also affects the fracture behaviour of annealed ferritic Ductile Iron. The influence of carbon content on notched impact properties is primarily on the upper shelf energy, which decreases with increasing carbon content, as shown in Figure 3.42. The influence of carbon in this region, in which fracture occurs by the formation of voids on graphite nodules, and the growth and coalescence of these voids, is to increase the number and size of nodules. Increasing carbon content thus reduces the plastic deformation required to grow and coalesce voids, resulting in reduced plastic fracture energy. This relationship between carbon content and limiting plastic fracture strain is consistent with the observation that elongation and other indicators of ductility in ferritic Ductile Iron increase with decreasing carbon content. (Fluidity, microstructural and shrinkage considerations normally require carbon levels above 3.2 per cent.)
The strong influence of silicon on the ductile-brittle transition temperature of ferritic Ductile Iron is shown in Figure 3.43. This Figure indicates that, to optimize low temperature toughness, silicon contents should be kept as low as possible. The successful production of as-cast carbide-free, low silicon Ductile Iron with a fully ferritic matrix requires high purity charge materials to minimize pearlite and carbide forming elements, controlled melting, holding and treating practices, and highly effective inoculation to maximize nodule count. The reduction in silicon level reduces both the yield and tensile strengths of the ferritic iron, and an offsetting addition of a less harmful ferrite strengthening element (such as nickel) is then needed to meet strength requirements. As with carbon, other considerations, especially microstructural control, require final silicon levels above 2 per cent.
Manganese, copper and nickel are some of the other major elements normally found in ferritic Ductile Iron. Manganese levels are kept low through dilution with high purity pig iron to avoid pearlite and carbide formation. The use of copper to strengthen low-silicon ferrite is precluded by its strong effect on transition temperature. A one per cent addition of copper will raise the transition temperature by 45oC (80oF). Nickel, which increases the transition temperature by only 10oC (20oF) for a 1 per cent addition, is the preferred ferrite strengthener for ferritic Ductile Irons requiring maximum low temperature toughness. Depending on its level, the pearlite stabilizing effect of the nickel may require an annealing treatment to ensure a fully ferritic matrix. Phosphorus, an impurity element in Ductile Iron, has a strong embrittling effect at levels as low as 0.02 per cent, see Figure 3.49.
Heat treatment, through its influence on microstructure, has a strong effect on impact properties. Figure 3.44 shows the effect on notched Charpy impact properties of the heat treatments described in Table 3.4. Subcritical annealing produced a fully ferritic, low strength structure with the highest upper shelf energy and the second lowest transition temperature. A special quench and temper treatment in which a low austenitizing temperature was used to produce a low carbon austenite, which was subsequently quenched and tempered, produced a superior combination of high strength and the low transition temperature. The normalized and tempered structure produced the poorest impact properties. When considering a material to obtain the best impact properties produced by the various heat treatments, it should be noted that the composition of the Ductile Iron used in the tests (shown at the bottom of Table 3.4) is "very poor". This material has high levels of phosphorus, silicon, chromium and manganese. The impace strength would be higher if the chemistry was improved.
*Analysis: 3.65% TC, 2.48% Si, 0.52% Mn, 0.65% P, 0.78% Ni, 0.08% Cr, 0.15% Cu
Table 3.4. Summary of heat treatments and tensile properties for the Ductile Iron samples used in Figure 3.44.
Impact properties of ferritic Ductile Irons are influenced by both nodularity and nodule count. In Figure 3.45, notched Charpy energies in the upper shelf region decrease significantly with decreasing nodularity. Transition temperatures and lower shelf energies are not affected by graphite shape. Nodule count also has a significant influence on both upper shelf energy and transition temperature. Increasing the nodule count from 180/mm2 to 310/mm2 (Figure 3.46) causes a decrease in transition temperature of 40oC (70oF) and a 25 per cent decrease in upper shelf energy. The use of late inoculation to produce higher and more consistent nodule counts presents both the designer and foundryman with a dilema. Should upper shelf energies be sacrificed in order to obtain increased low temperature impact properties? The Charpy test is too imprecise and potentially erroneous to answer this question. Fracture mechanics may be required to determine the true contribution of nodule count to fracture toughness. Figure 3.48 suggests that increasing the nodule count from a low level, may improve fracture toughness when Ductile Irons exhibit brittle fracture behaviour.
Although tensile data indicate that Ductile Iron has strength and ductility similar to cast steels, standard notched Charpy tests suggest that Ductile Iron has significantly lower fracture toughness, with energy values in the range 16-24 joules (12-18 ft lbf) compared to cast steels with 60-75 joules (44-55 ft lbf). Before Charpy data is used to disqualify Ductile Iron from critical applications because of its apparently inferior toughness, the following shortcomings of the Charpy test should be considered in the light of current fracture mechanics information to determine toughness. First, fracture mechanics samples are precracked, while Charpy notches are relatively blunt. As a result, fracture mechanics tests measure resistance to crack propagation, while Charpy tests measure both initiation and propagation. Second, fracture toughness tests are conducted under quasi-static stress conditions while the Charpy test involves impact loading. Finally, fracture mechanics test samples are large enough to produce plane strain conditions, while the Charpy test involves plane stress, a fact clearly confirmed by the shear lips on fractured steel Charpy samples tested in the upper shelf region.
The formation of shear lips is the underlying cause of the significant difference between the Charpy behaviour of Ductile Iron and cast steel. The shear lips developed by the steel are responsible for a considerable fraction of its upper shelf energy. Due to the strain-limiting nature of the coalescence of voids initiated on graphite nodules, Ductile Iron does not exhibit shear lip formation under any conditions. As a result, when tested under the "similar" plane stress conditions present in the Charpy test, the shear lip formation of steel produces a significantly higher upper shelf fracture energy than Ductile Iron. Under plane strain conditions that could be expected in many component failures, the "shear lip advantage" of steel would be absent, with dramatically lower fracture toughness.
To eliminate the differences in upper shelf fracture mode between cast steel and ferritic Ductile Iron, the Charpy test was modified, using precracked and side- grooved samples to provide plane strain conditions at the initiation of crack growth. Using the J-integral method, the dynamic stress intensity factor KID was calculated for both materials over a temperature range including both brittle and ductile fracture modes. Figure 3.47 shows that the fracture toughness of cast steel was superior to that of ferritic Ductile Iron at temperatures above 90oF (32 oC) but that the superiority was much less than that suggested by the Charpy test. Due to a much lower ductile-to-brittle transition temperature, Ductile Iron exhibited superior fracture toughness below 90o F (32o C).
Figure 3.47 indicates that the fracture toughness of good quality ferritic Ductile Iron is excellent to temperatures as low as -80oF (-62oC), giving a KID of 37.5 ksi (square root) in. (41 MPa (square root) m), which corresponds to a critical flaw size of 0.5 in. (1.25 cm) for a design stress equal to the yield stress, applied under static fracture conditions. Above 0 oF (- 18oC), the KID is 80 ksi (square root)in. (87 MPa (square root) m) giving a critical flaw size of 1.5 in. (3.75 cm). Both flaw sizes can be detected and prevented by the quality assurance and production procedures practiced by competent Ductile Iron foundries. Assuming such flaws can be avoided, ferritic Ductile Iron can be considered sufficiently tough to resist unstable crack propagation at temperatures as low as -80o F (-62oC).
Figure 3.48 illustrates the relationship between fracture toughness and nodule count for pearlitic Ductile Iron tested at room temperature. This level of fracture toughness, at a temperature well below the transition temperature for pearlitic irons, (see Figure 3.41) indicates that these irons are tougher than indicated by the notched Charpy test and have good flaw tolerance at temperatures at which they are labeled "brittle" by the Charpy test. The relationship between fracture toughness and modularity indicates that the nodules are playing a role in determining fracture toughness, possibly through the relaxation of triaxial stresses through void formation at the crack tip.
The dynamic tear test, an accepted ASTM fracture test method, overcomes many of the shortcomings of the Charpy test and is cheaper and more suitable than plane strain fracture toughness testing for production testing of ferrous castings. This test has become widely accepted in the automotive industry and has been made mandatory for the characterization of the fracture properties of castings used in critical applications. To ensure validity of test results, the dynamic tear specimens are cast to size in the foundry and tested full size to replicate performance of an actual casting. Figures 3.49, 3.50 and 3.51 illustrate dynamic tear behaviour for as-cast ferritic and pearlitic, and annealed ferritic Ductile Irons respectively. When compared to similar Charpy data in Figure 3.41, the dynamic tear data in Figures 3.50 reveals a slightly higher transition temperature for the ferritic sample but significantly lower transition temperatures for the pearlitic grades. Figure 3.49 is noteworthy for two features: the low transition temperature of the low phosphorus, annealed ferritic iron, and the significant increase in transition temperature and reduction in upper shelf energy produced by an increase in phosphorus content to 0.05 per cent.
Figure 3.51 compares the dynamic tear data for four ferritic cast irons. A full, ferritizing anneal reduces the fracture transition temperature and increases the upper shelf energy of Ductile Iron, compared to an as-cast ferritic structure. The use of a subcritical anneal instead of a normal full ferritizing treatment resulted in a similar transition temperature but a higher upper shelf energy. In addition to slightly better impact properties, a subcritical anneal also produces improved fatigue strength.
Temper embrittlement, as found in certain quenched and tempered steels, may also occur in similarly treated Ductile Irons with susceptible compositions. This form of embrittlement, which does not affect normal tensile properties but causes significant reductions in fracture toughness, can occur in Ductile Irons containing high levels of silicon and phosphorus which have been tempered in the range 650-1100oF (350-600oC) and cooled slowly after tempering. Although normally associated with tempered martensitic matrices, temper embrittlement can also occur if the matrix is tempered to the fully ferritic condition. Temper embrittlement can be prevented by keeping silicon and phosphorus levels as low, adding up to 0. 15 per cent molybdenum and avoiding the embrittling heat treating conditions.
This form of embrittlement, named because it may be found in certain galvanized Ductile Iron and Malleable Iron castings, does not involve zinc and the galvanizing process directly but is caused in castings with relatively high silicon and phosphorus levels by the thermal environment created during galvanizing. For example, an annealed ferritic Ductile Iron of susceptible composition will be embrittled by quenching or rapid cooling after galvanizing in the temperature range 650-950oF (350-500oC). Although galvanizing embrittlement and temper embrittlement are both related to high silicon and phosphorus levels, they differ in other, important respects. Galvanizing embrittlement normally occurs in annealed ferritic castings and is caused by rapid cooling from the embrittling temperature range, while temper embrittlement occurs in quenched and tempered castings and is caused by slow cooling.
The Modulus of Rigidity, or Modulus of Elasticity in Torsion, is the ratio of shear stress to shear strain. The Modulus of Rigidity and Modulus of Elasticity are related by the equation:
E = 2G(1 + v)= 2.55G
= Modulus of Elasticity,
The 0.2% compressive yield strength can be up to 20% higher than the tensile yield strength measured at the same offset. The relationship between the compressive yield strength and Brinell hardness is shown in Figure 3.52. The proportional limit in compression is a slightly higher proportion of the compressive yield strength and does not vary significantly between grades. A suitable estimate of the proportional limit in compression is obtained by using 75% of the 0.2% compressive yield strength for all grades of conventional Ductile Iron.
The ultimate strength in shear or torsion is generally considered to be about 90% of the tensile strength. However, there is a scarcity of accurate shear strength values in materials such as Ductile Iron that show some ductility because in a double shear test it is very difficult to avoid bending. Data on the proportional limit and yield stress in torsion are more reliable, with torsional values being about 75% of the respective tensile values.
Damping capacity plays a significant role in modern engineering design. High damping capacity reduces the noise and subsonic vibrations emitted by machinery components which are subjected to cyclic stressing. Combustion engines are quieter and transmit less vibration to attached components, machine tools are less noisy and produce a smoother surface finish. The only disadvantages of damping are the additional frictional losses and related heat build-up that result from the absorption of vibrational energy by a material with high damping capacity.
A more profound significance of damping capacity is in its contribution to fatigue resistance. Two materials with the same measured fatigue resistance but with different damping capacities will perform differently in service. In most actual service environments (as opposed to fatigue testing) vibrations occur intermittently and with varying frequency and amplitude. A rapid damping of these vibrations reduces the length of time during which the stress amplitude may reach or exceed the fatigue limit. As a result, high damping capacity enhances fatigue resistance.
Damping is the ability of a material to absorb vibrational energy by some form of internal friction. In metals the primary damping mechanism is localized non-elastic (microplastic) behaviour. Under cyclic loading conditions this microplastic behaviour, shown in Figure 3.53, produces a hysteresis loop whose area is proportional to the energy absorbed during each cycle (vibration). The low stress behaviours of Gray Iron, Ductile Iron and mild steel, see Figure 3.3, indicate their relative damping capacities. Gray Iron, which exhibits non-elastic behaviour at very low stresses, has the highest damping capacity, while steel, which behaves elastically up to its yield point, has the lowest damping capacity. Figure 3.54 schematically illustrates the relative damping capacities of these materials through a comparison of reduction in vibrational amplitude with time. The relative decreases in vibrational amplitude illustrated in this Figure can vary as follows for ferrous materials:
In addition to the general variations related to different types of material, damping capacity is also affected within a family of materials by applied stress state and microstructure. Figure 3.55 shows the variation in damping capacity with surface stress for Gray Iron, low carbon steel and ferritic and pearlitic Ductile Iron. Figures 3.56 and 3.57 illustrate the influence of microstructure on damping behaviour of Ductile Iron. As would be expected from the relative damping capacities of Gray and Ductile Irons, as the percentage of spherical graphite decreases ( and the amount of flake-like graphite increases), damping capacity increases significantly (see Figures 3.56). This Figure also shows that damping capacity is not affected strongly by carbide contents up to 14 volume per cent. Figure 3.57 shows that damping capacity generally decreases with increased matrix hardness and increases with carbon content. The only exception to the damping-hardness relationship is for as-quenched martensite, in which the internal stresses produced by the formation of martensite increase microplastic deformation and thus increase damping. As shown in Figure 3.55, ferritic and pearlitic Ductile Irons exhibit a transition in relative damping capacity as the applied stress is increased. At low stresses, the softer ferritic matrix has higher damping capacity, while at higher stresses, the damping capacity of the pearlitic matrix is greater.
The generally accepted value for the room temperature density of Ductile Iron is 7.1 g/CM3. Density is affected primarily by the percentage of graphitized carbon, with densities varying from 6.8 g/CM3 to 7.4 g/CM3 for high carbon ferritic and low carbon pearlitic irons respectively. The density of a typical cast steel - 7.8 grams/CM3 - is almost 10 per cent higher than that of Ductile Iron. The replacement of a steel casting or forging with a lighter Ductile Iron casting improves the component strength: weight ratio, reducing energy savings and lifetime costs, especially in reciprocating components such as automotive crankshafts.
The coefficient of linear thermal expansion of Ductile Iron depends primarily on microstructure, although it is influenced to a minor extent by temperature and graphite structure. In unalloyed Ductile Iron, composition has only a slight influence on thermal expansion, but alloyed austenitic Ductile Irons can exhibit significantly different expansion behaviour, see Table 3.5.
Table 3.5. Effect of temperature on the coefficient of thermal expansion for different Ductile Irons.
The thermal and electrical conductivity of Gray and Ductile Irons are influenced strongly by graphite morphology. The conductivity is higher in Gray Iron because of the semi-continuous nature of the graphite flakes. Because of the influence of flake graphite on the conductivity, the volume fraction of graphite plays an important role in Gray Iron, but not in Ductile Iron. In addition to graphite shape, microstructure, composition, and temperature also influence thermal conductivity. Ferritic Ductile Irons have a higher thermal conductivity than pearlitic grades, and quenched and tempered irons have values between those of ferritic and pearlitic irons. In the range 20-500oC (68-930oF), the thermal conductivity of ferritic grades is 36 W/m oK (250 Btu in. / ft2 hoF). Conductivity for pearlitic grades over the same temperature range is approximately 20 per cent less.
Specific heat, the amount of energy required to increase the temperature of a unit mass of a body by one degree, generally increases with temperature, reaching a maximum whenever a phase transformation occurs. For unalloyed Ductile Iron. the specific heat varies with temperature as follows:
Ductile Irons, with discontinuous spherical graphite, have lower electrical resistivity than Gray Irons which have semi-continuous flake graphite. The primary elements effecting resistivity are silicon and nickel, both of which increase resistivity. The effects of matrix structure and silicon content on the electrical resistivity of Ductile Iron at room temperature are shown in Figure 3.58.
The magnetic properties of Ductile Irons are determined mainly by their microstructures. The spheroidal shape of the graphite particles in Ductile Irons gives them higher induction and higher permeability than Gray Irons with a similar matrix. Ferritic Ductile Irons are magnetically softer than pearlitic grades - they have higher permeability and lower hysteresis loss. For maximum permeability and minimum hysteresis loss, ferritic, low phosphorus irons should be used. Magnetization and permeability curves are shown in Figures 3.59 and 3.60 for three ferritic cast irons. The magnetic and electrical properties of these irons are summarized in Table 3.6.
Table 3.6 Magnetic
and electrical properties of the ferritic irons
Mechanical wear may be defined as surface deterioration and/or material loss caused by stresses arising from contact between the surfaces of two bodies. Wear is primarily mechanical in nature but chemical reactions may also be involved. Wear is a complex phenomenon and may involve one or more of the following mechanisms:
The complexity of wear phenomena and their dependence on both material properties and environment have precluded the use of a universal wear test to evaluate and compare the wear behaviour of different materials under different wear conditions. As a result, many tests have been developed for evaluating wear resistance, with each test applying to a specific set of conditions. Therefore, the discussion of wear resistance is limited to general, comparative statements and to some thoughts on how the microstructure of Ductile Irons affect their wear resistance.
Cast irons have been recognized for many years as ideal materials for a wide range of wear applications, especially frictional wear under both dry and lubricated conditions. In dry wear, sufficient heat may be generated by the friction between the working surfaces to harden the individual surfaces or, in severe cases, fuse them together. Under these conditions the graphite particles in cast irons lubricate the surfaces, reducing friction and minimizing surface deterioration by overheating. Graphitic cast irons also perform well in lubricated sliding wear. The graphite particles on the wear surfaces act as reservoirs for oil and, under loads high enough to displace the oil film, the lubricating effect of the graphite itself provides galling resistance.
The wear resistance of Ductile Irons are determined primarily by their microstructures. The presence of 8-11 volume per cent graphite provides both the graphitic lubrication and oil retention essential to some wear applications. Pearlite, consisting of very hard lamellar carbide in a soft, ductile matrix of ferrite, exhibits good wear resistance under wear conditions involving both friction and moderate abrasion. Further improvements in resistance to abrasive wear may be obtained through alloying and/or heat treatment to produce a harder martensitic, austempered or bainitic matrix, Figure 3.61. Additional information on heat treatment and surface treatment can be found in Sections VII and Section IX.
Unalloyed Ductile Irons exhibit approximately the same corrosion resistance as Gray Iron and are superior to unalloyed steel, and even highly alloyed steels in certain environments. Corrosive environments degrade the performance of Ductile Iron in two ways: the embrittlement of monotonically stressed components described earlier in this section, and the loss of material and structural integrity caused by corrosive action alone. Corrosion can also play a significant role in abrasive wear resistance. Corrosion of Ductile Irons and other ferrous materials is a complex phenomenon and a detailed discussion of corrosion behaviour is beyond the scope of this section. The corrosion behaviour of alloyed Ductile Irons is discussed briefly in Section V. Data describing the general corrosion behaviour of Ductile Irons can be found in the following sources:
"Corrosion Data Survey", National Association of Corrosion Engineers, Katy, Texas, 1974.
"Iron Castings Handbook", Iron Castings Society, 1981.
A Design Engineer's Digest of Ductile Iron, 5th Edition, 1983, QIT-Fer et Titane Inc., Montreal, Quebec, Canada.
BCIRA Broadsheet 157-2, "Stress/strain behaviour of nodular and malleable cast irons," British Cast Iron Research Association, 1981.
B. V. Kovacs, "Quality Control and Assurance by Sonic Resonance in Ductile Iron Castings," Transactions, American Foundrymen's Society, Vol. 85, 1977, pp. 499-508.
B. V. Kovacs, "On the Interaction of Acoustic Waves with SG Iron Castings." Transactions, American Foundrymen's Society, Vol. 83, 1975, pp. 497-502.
W. Siefer and K. Orths, "Evaluation of Ductile Iron in Terms of Feasible Properties of the Material." Transactions, American Foundrymen's Society, Vol. 78, 1970 pp. 382-387.
D. L. Crews, "Quality and Specifications of Ductile Iron." Transactions, American Foundrymen's Society, Vol. 82, 1974, pp. 223-228.
D. Venugopalan and A. Alagarsamy, "Effects of Alloy Additions on the Microstructure and Mechanical Properties of Commercial Ductile Iron." Transactions, American Foundrymen's Society, Vol. 98, 1990, Paper #90-122.
E. N. Pan, M. S. Lou and C. R. Loper Jr., "Effects of Copper, Tin and Manganese on the Eutectoid Transformation of Graphitic Cast Irons." Transactions. American Foundrymen's Society, Vol. 95, 1987, pp. 819-840.
T. Levin, P.C. Rosenthal, C.R. Loper Jr. and R. Heine, "Tin and Copper in Ductile Iron." Transactions, American Foundrymen's Society, Vol. 79, 1971, pp. 493-514.
0. Okarafar and C. R. Loper, "The Effect of the Exposure to Water on the Tensile Properties of Ductile Cast Iron.", Transactions, American Foundrymen's Society, Vol. 87, 1979.
G. J. Cox, "The impact properties of different types of ferritic spheroidal graphite cast iron," British Foundryman, 1973.
K. B. Palmer, "The mechanical and physical properties of engineering grades of cast iron at subzero temperatures," British Cast Iron Research Association, Oct., 1988.
P. J. Rickards, "The low-temperature tensile properties of nodular graphite cast irons," Cryogenics, Dec., 1969.
C. R. Wilks, N. A. Mathews, and R. W. Kraft, Jr., "Elevated Temperature Properties of Ductile Cast Irons," Transactions, American Society for Metals, 1954.
F. B. Foley, "Mechanical Properties at Elevated Temperatures of Ductile Cast Iron," Transactions, American Society for Metals, 1956.
R. O. Schelling and J. T. Eash, "Effect of Composition on the Elevated Temperature Properties of Ductile Iron," Proceedings, American Society for Testing and Materials, 1957.
K. B. Palmer, "High Temperature Properties of Cast Irons," Engineering Properties and Performance of Modem Iron Castings, British Cast Iron Research Association, Birmingham, England, 1970.
D. L. Sponseller, W. G. Scholz, and D. F. Rundle, "Low Alloy Ductile Irons for Service at 1200-1500o F," Transactions, American Foundrymen's Society, 1968.
Miller and Company, The Ductile Iron Process, 1972.
The Iron Castings Handbook, Iron Castings Society, Inc., 1981.
C. Isleib and R. Savage, "Ductile Iron-Alloyed and Normalized," Transactions, American Foundrymen's Society, Vol. 65, 1957.
S. Palmer, "What is Required of Iron Castings As Engineering Components?" Engineering Properties and Performance of Modern Iron Castings, British Cast Iron Research Association, Birmingham, England, 1970.
S. I. Karsay, Ductile Iron II, Quebec Iron and Titanium Corporation, 1972.
G. N. J. Gilbert, Tensile and Fatigue Tests on Normalized Pearlitic Nodular Irons," journal of Research, British Cast Iron Research Association, Vol. 6, No. 10, February 1957 pp. 498-504.
K. B. Palmer and G. N. J. Gilbert,
"The Fatigue Properties of Nodular Cast Iron, " journal
R. C. Haverstraw and J. F. Wallace,
"Fatigue Properties of Ductile Iron," Gray & Ductile
A. G. Fuller, "Effect of Graphite Form on Fatigue Properties of Pearlitic Ductile Irons." Transactions, American Foundrymen's Society, Vol. 85, 1977, pp. 527-536.
M. Sofue, S. Okada and T. Sasaki, "High-Quality Ductile Cast Iron with Improved Fatigue Strength." Transactions, American Foundrymen's Society, Vol. 86, 1978, pp. 173-182.
K. B. Palmer, "Fatigue Properties of Cast Irons, " Engineering Properties and Performance of Modern Iron Castings, British Cast Iron Research Association, Birmingham, England, 1970.
J. F. Janowak, A. Alagarsamy, and D. Venugopalan, "Fatigue Strength of Commercial Ductile Irons. ", Transactions, American Foundrymen's Society, Vol. 98, 1990, Paper #90-123.
Int. Komitee Giessereitechn, Vereinigungen, Fatigue Strength of Nodular Iron, VDG Technical Report No. 017, Dusseldorf, 1975.
Kolene Corporation, Detroit Michigan, Chart 6032 is Figure 3.36
K. Roehrig, "Thermal Fatigue of Gray and Ductile Irons.", Transactions, American Foundrymen's Society, Vol. 86, 1978, pp. 75-88.
G. Sandoz, H. Bishop and W. Pellini, "Notch Ductility of Nodular Irons," Transactions, American Foundrymen's Society, 1954.
H. E. Mead, Jr. and W. L. Bradley, Fracture Toughness Studies of Ductile Cast Iron Using a J-Integral Approach, Transactions, American Foundrymen's Society, 1985.
K. E. Kinney, W. L. Bradley and P. C. Gerhardt Jr., "An evaluation of the Toughness of Ductile Iron vs. Cast Steel Using Modified Charpy Test Specimens. " Transactions, American Foundrymen's Society, Vol. 92, 1984, pp. 239-250.
P. F. Timmins, "The Effect of Microstructure on Defect Tolerance in Pearlitic Spheroidal Graphite Iron." private communication, courtesy of Highland Foundry Ltd., Surrey, British Columbia, Canada.
BCIRA Broadsheet 212, "Factors influencing the ductile or brittle behaviour of nodular irons," British Cast Iron Research Association, 1982.
G. N. J. Gilbert, Engineering Data on Nodular Cast Irons, British Cast Iron Research Association, Birmingham, England, 1974.
Ductile Iron, Metals Handbook, American Society for Metals, 9th edition, Vol. 15, 1988.
L. Jenkins, "Ductile Iron - an Engineering Asset.", Proceedings of the First International Conference on Austempered Ductile Iron, American Society for Metals, 1984.
R. Barton, "Embrittlement of nodular (SG) iron: its forms and prevention," Fourth International Conference of Licensees for the + GF + Converter Process, Schaffhausen, Switzerland, 1981.
G. Dinges and W. Gruver, "Temper Embrittlement of Ductile Iron," Transactions, American Foundrymen's Society, 1966.
R. A. Flinn, P. K. Trojan, and D. J. Reese, "The Behavior of Ductile Iron Under Compressive Stress," Transactions, American Foundrymen's Society, 1960.
BCIRA Broadsheet 203-5, "Density of cast iron." British Cast Iron Research Association. 1984.
Ductile Iron, International Nickel Company, New York, New York, 1956.
A. B. Everest, "Spheroidal-Graphite Cast Iron," Foundry Trade journal, 1950, pp. 57-64.
BCIRA Broadsheet 203, "Thermal conductivity of unalloyed cast irons," British Cast Iron Research Association, 1981.
L. W. L. Smith, "The determination of the thermal conductivity of cast iron." BCIRA journal, Vol. 21, 1973.
BCIRA Broadsheet 203-3, "Electrical resistivity of unalloyed cast iron," British Cast Iron Research Association, 1984.
K. B. Palmer, "The electrical resistivity of cast iron," BCIRA journal, June, 1953.
M. Decrop, "Le Magnetisme des Fontes," Fonderie, July, 1963.
BCIRA Broadsheet 203-7. "Specific heat capacity of cast irons." British Cast Iron Research Association, 1985.
H. T. Angus, Cast Iron: Physical and Engineering Properties, 2nd ad., Butterworths Inc., 1967.
C. F. Walton, Gray Iron castings Handbook, Gray and Ductile Iron Founders' Society, Cleveland, Ohio, 1958.
BCIRA Broadsheet 203-5, "Density of cast irons, " British Cast Iron Research Association, 1984.
G. N. J. Gilbert and C. L. Pidgeon, "The growth and scaling properties of cast irons in air at temperatures of 350oC and 400oC for times up to 21 years," BCIRA Journal, July, 1982.
E. Plenard, "Cast Iron Damping Capacity, Structure and Related Properties," Transactions, American Foundrymen's Society, Vol. 70, 1962, pp. 298-305.
E. Plenard, Damping Capacity of Cast Iron, Foundry Trade journal, Oct, 1966, pp. 541-549.
Annual Book of ASTM Standards, Volume 01.02, Ferrous Castings, 1987
Cast Metals Handbook, American Foundrymen's Society, 1957.