High Temperature Thermal Fatigue Property of Thin-Section Ductile Cast Iron · · 2013-11-25Ductile Cast Iron E. N. Pan, C. C. Fan, ... both the nodule count and nodularity increase

High Temperature Thermal Fatigue Property of Thin-Section Ductile Cast Iron

E. N. Pan, C. C. Fan, H. Y. Chang

National Taiwan University, Taipei, Taiwan

Copyright 2010 American Foundry Society

ABSTRACT

The purpose of this research is to establish the optimal casting conditions for producing thin-section (2-3 mm) spheroidal graphite cast irons for thermal cyclic applications (up to 800C [1472F]). Experimentally, the microstructures and the constrained thermal fatigue property will be evaluated and correlated with the alloy design and casting parameters. The results show that, for a fixed C content of some 3.0%, both the nodule count and nodularity increase first with increasing Si content, reach maxima at around 3.8-3.9%Si (4.27-4.30%CE), and then decrease with further increase in Si content. On the other hand, %ferrite increases gradually with increasing Si content, reaches a maximum at around 4.6-4.8%Si (4.53-4.60%CE), and then remains more or less constant or decreases slightly. Furthermore, irons with 4.45-4.55%CE exhibit somewhat higher nodule count and %ferrite than irons with 4.75-4.85%CE. The optimal alloy design for attaining the best thermal fatigue property has been found to be a moderate hyper-eutectic composition, i.e., 4.5-4.6%CE, with a combination of relatively low C content and high Si content, e.g., 3.0%C + (4.6-4.8)%Si. In addition, the best thermal fatigue property corresponds to a microstructure with the highest ferrite content and moderate nodule count. Furthermore, adding some 0.5%Mo to the iron significantly increases the thermal fatigue life. INTRODUCTION Development of lighter weight parts for vehicle components to meet the challenges to reduce emissions and improve fuel economy has been a continuing task in the automotive industry. Vehicle weight reduction propels the substitution of aluminum or magnesium for cast iron and steel in a number of automotive components, due to its lower density compared to cast iron and steel. The use of aluminum or magnesium, however, results in increased vehicle costs. Since ductile iron has a comparatively lower cost, thin-walled castings should gain an edge on the stiff competition among materials chosen for automotive applications. Thin-walled ductile iron castings (3mm or less) not only satisfy the requirement for lower weight, but also offer the design engineer higher strength to weight ratios and better performance for high temperature applications than many alternative materials.1,2

The production of relatively thin section ductile iron castings may encounter the problem of carbide formation because of the extremely high solidification cooling rates. Nevertheless, this difficulty can be overcome by choosing appropriate inoculants and/or employing effective inoculation methods, such as late or stream inoculation and in-mold process., Therefore, carbide-free thin-walled ductile iron castings can be obtained. Inoculation not only improves graphite nodularity and increases nodule count, but also prevents carbides from forming by providing new or additional nuclei on which solidification and growth of the stable graphite-austenite eutectic can occur. Currently, several proprietary inoculants that contain various minor active inoculating elements, e.g., Al, Ca, Ba, Zr and Ce, are available. The amount of inoculant used varies depending on the section size of the castings being produced, the effectiveness of the inoculant and also the method of addition. In this study, a proprietary Fe-Si inoculant that bears small amounts of Al, Ca and Ba was used. In the past, a great deal of efforts have been made to establish the production technologies of thin-walled ductile iron castings, including control of microstructure, solidification modeling, dimensional capability and generation of useful mechanical property data.3-10 The thin-walled ductile iron castings for automotive applications often required good high temperature performance, especially cyclic thermal fatigue. Regarding the evaluation of thermal fatigue property, Park et al.11,12 employed a constrained thermal fatigue test apparatus, where the test specimen was heated by high frequency induction and cooled by conduction of heat to the water-cooled grips. The test specimens used are hollow cylindrical shape. In this study, a constrained thermal fatigue testing device, similar to Lee et al.’s13, was employed for thin plate type test specimens (the details are described under Experimental Procedures). It has been well realized that thermal fatigue property is influenced mainly by the tensile strength of the alloy and also the resistance to high temperature stress relaxation. Both properties in turns are related with the alloy composition and microstructure of the alloys tested.14-20

Concerning the alloy composition, the contents of C and Si (or carbon equivalent, CE) affect the degree of chilling tendency, carbon flotation, solidification shrinkage and castability. For castings with relatively thin sections, somewhat higher CE values are required to

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prevent carbides formation, but the CE shouldn’t be too high to prevent the occurrence of carbon flotation and misrun. For section thickness less than 5mm, CE of 4.5-4.7% is recommended. Furthermore, Si is of great importance in high temperature applications. Increasing Si content raises the critical temperature, allowing the as cast phases to remain intact until a higher temperature. Also, the formation of a thicker SiO2 oxide film on the casting surface due to an increased Si content improves oxidation resistance. However, too high a Si content may cause room temperature brittleness. Therefore, the optimal Si range of 4.0-6.0%Si is recommended for high temperature applications. In addition, Mo stabilizes carbides (or cementites in pearlite), resists high temperature stress relaxation and creep. As suggested, the addition of Mo is less than 2.0%, with the optimal addition of 0.4-0.6%.11,12,14,15 In this study, the CE value of 4.3-4.8%, the Si content of 3.0-5.0%, and also the effect of 0.5%Mo addition were investigated. Therefore, this study was conducted in anticipation of gaining better understanding of how the alloy design and microstructure affect the thermal fatigue property of thin-walled ductile iron castings. EXPERIMENTAL PROCEDURES OBJECTIVE The purposes of this research are twofold: 1) to study the influence of chemical composition (C, Si, CE, Mo) and processing parameters (section thickness, molding material and pouring temperature) on microstructures of thin-section ductile cast irons; and 2) to correlate the alloy design and microstructure with thermal fatigue life of thin-section ductile cast irons. ALLOY DESIGN To assess the effects of C, Si (or carbon equivalent, CE) and Mo on the microstructure and thermal fatigue property of thin-section ductile cast irons, three different carbon equivalent (CE) values were examined, i.e., 4.40%CE, 4.55%CE and 4.70%CE. For each carbon equivalent, three nominal levels of Si were considered, i.e., 3.0%, 4.0% and 5.0%, with the C content being adjusted accordingly to meet the designed carbon equivalent value. In addition, 0.5%Mo was added in one specific heat to evaluate the effect of Mo addition on thermal fatigue life. The alloy design is given in Table 1, whereas the final chemical analyses for all heats are presented in Table 2. METALLURGICAL (PROCESS) PARAMETER The process parameters to be investigated include section size, molding medium and pouring temperature. For each heat, plate castings with two different section thicknesses, 2 mm and 3 mm, were cast at two different temperatures, 1450±15C (2642±27F) and 1400±15C (2552±27F), into two different molding media, green sand and chemically- bonded sand. The geometry and dimensions of patterns for test castings are indicated in Fig. 1. A series of 3

letters and 1 number were used for specimen codes, where the first 3 letters denote Heat number (A–K), level of pouring temperature (H: high pouring temperature, 1450±15C (2642±27F), L: low pouring temperature, 1400±15C (2552±27F) and type of molding medium (G: green sand, C: chemically-bonded sand), respectively, and the last number indicates section thickness (2: 2 mm; 3: 3 mm). For instance, specimen A-HC2 denotes Heat A, poured at a higher temperature of 1450±15C (2642±27F) and cast into a chemically-bonded sand mold of 2 mm section. Fig. 1. Dimensions of patterns used in this study, (a) 2mm-plate pattern:2×60×250 mm, (b) 3mm-plate pattern: 3×60×300 mm. MELT PREPARATION A total of 11 heats were prepared in a 50 kW, high- frequency (3000Hz) induction furnace using charges composed of low sulfur pig iron, steel scrap and Fe-Si alloy (for Si content adjustment). The spheroidization treatment was carried out at about 1500C (2732F) by adding 1.2% of an Mg-FeSi alloy containing minor amounts of rare earth (RE), Ca and Al using the sandwich method, where the nodularizer was sealed in a steel-can and placed at the bottom of the treatment ladle. Inoculation was performed by stirring the inoculant (0.6%) into the melt stream while the melt was transferred into a ladle for spheroidization. The inoculant employed is a proprietary Fe-Si alloy that contains Ba, Ca and Al. The chemical compositions of charge materials and treatment alloys employed in this study are presented in Table 3.

(a)

(b)


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Table 1. Alloy Design (mass %)

Nominal compositions

C.E.* C Si Mo

3.40 3.0 3.06 4.0 4.40 2.73 5.0 3.55 3.0 3.21 4.0 2.88 5.0

---

4.55

2.88 5.0 0.5 3.70 3.0 3.36 4.0 4.70 3.03 5.0

---

＊ C.E.=C%+1/3Si% MICROSTRUCTURE OBSERVATIONS Microstructure analyses, including nodule count, graphite nodularity and matrix structure, were performed on samples taken from the center of both 2 mm and 3 mm section plate castings. The nodule count, graphite nodularity and matrix structure were measured by a commercial image analyzer. The measurements of nodule count were performed under the image magnification of 100 times. The selected minimum diameter of graphite nodules was 5μm. THERMAL FATIGUE TEST The experiment apparatus, both schematic and the actual set-up, for evaluating the thermal fatigue life is illustrated in Fig. 2. The thermal fatigue test specimen was mounted at both ends on a fixture by a fixed clamping force of 5 ft-lb. An induction coil placed over the test specimen was used for heating. The thermal fatigue specimens include both 2 mm and 3 mm sections with a V-notch in the middle. The depths of the V-notch are 0.25 mm and 0.50 mm for 2 mm and 3 mm sections, respectively, as shown in Fig. 3. During the cyclic thermal fatigue testing, firstly, power was turned on and the specimen was heated to the pre-determined temperature, 800C (1472F), in 60 seconds, and then the power was turn off and the test specimen

was immediately quenched to room temperature by spraying water directly onto the test specimen. The cooled test specimen was dried by compressed air, and then the test specimen was ready for the next thermal cycle. The procedures continued until the test specimen failed completely. Fig. 2. The schematic drawing and actual set-up of the thermal fatigue test apparatus.

Table 3. Chemical Compositions of the Charge Materials and Treatment Alloys (mass %)

Size C Si Mn S P Mg RE Ca Al Ba Fe

QIT Pig Iron -- 4.33 0.13 0.022 0.014 0.027 -- -- -- -- -- Bal

Steel Scarp -- 0.2 0.3 -- -- -- -- -- -- -- -- Bal

Ferro-silicon -- -- 75.0 -- -- -- -- -- -- -- -- Bal

Nodularizer 0.5-5mm -- 42.0 -- -- -- 6.25 2.0 2.0 0.68 -- Bal

Inoculant 1-3mm -- 72.5 -- -- -- -- -- 2.25 1.0 2.0 Bal

Table 2. Chemical Analyses of All Heats (mass % )

Heat CE C Si Mn P S Mg Mo Fe A 4.00 3.01 2.98 0.04 0.03 0.01 0.015 --- Bal. B 4.34 3.28 3.19 0.04 0.03 0.01 0.019 --- Bal. C 4.15 2.88 3.82 0.05 0.02 0.01 0.025 --- Bal. D 4.25 2.58 5.00 0.04 0.02 0.01 0.032 --- Bal. E 4.46 3.42 3.09 0.03 0.03 0.01 0.013 --- Bal. F 4.56 3.15 4.22 0.04 0.03 0.01 0.019 --- Bal. G 4.50 2.91 4.77 0.07 0.02 0.01 0.036 --- Bal. H 4.84 3.85 2.96 0.03 0.03 0.01 0.029 --- Bal. I 4.74 3.34 4.20 0.06 0.02 0.01 0.064 --- Bal. J 4.76 3.06 5.11 0.05 0.02 0.01 0.016 --- Bal. K 4.61 2.76 5.54 0.08 0.03 0.01 0.026 0.53 Bal.


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Fig. 3. Dimensions of thermal fatigue test specimens. RESULTS AND DISCUSSION Effects of section size and some selected metallurgical and process parameters, including chemical composition, molding material and pouring temperature, on the microstructure and thermal fatigue life of relatively thin-section ductile irons were evaluated in this study. The thermal cyclic temperature range was between room temperature and 800C (1472F). THE EFFECT OF Si CONTENT AT A FIXED C CONTENT OF 3.0% ON MICROSTRUCTURE The effect of Si content on graphite morphology for both 2 mm- and 3 mm-section ductile irons cast in chemically-bonded sand molds is shown in Fig. 4. As shown, both nodule count and graphite nodularity increase first with increasing Si content, reach maxima at around 3.8-3.9%Si (1318 #/mm2, 94% for 2 mm-section and 1154 #/mm2, 96% for 3 mm-section), and then decrease with a further increase in Si content. In addition, the nodule count of 2 mm sections is higher than that of 3 mm sections by some 16%, while no clear difference in graphite nodularity can be observed between both sections. Similar results were also obtained for ductile irons cast in green sand molds. As for the matrix structure, changes in %ferrite (also %pearlite) as a function of %Si in both 2 mm and 3 mm section castings were plotted in Fig. 5 for green sand castings. As shown, the %ferrite increases gradually with increasing %Si, reaches maxima at around 4.6-4.8%Si, and then decreases slightly with a further increase in %Si. Also, the %ferrite in 3 mm sections is higher than that in 2 mm sections by 5% due to a lower solidification cooling rate in the former than in the latter. Fig. 6 illustrates a series of microstructures, which show that the highest nodule count was obtained in Heat C (3.8%Si, 1318 #/mm2), Fig. 6c and d, while the highest ferrite content occurred in Heat G (4.7%Si, 89% & 93%), Fig. 6e and f. Regarding the effects of casting parameters on microstructure, the results show that higher nodule counts can be obtained in castings with a thinner section or with a lower pouring temperature. However, the effect of molding material on nodule count is not significant. On

the other hand, all the afore-mentioned three casting parameters exert little influence on nodularity. Furthermore, the ferrite percentage is higher in castings with a thicker section, a higher pouring temperature and molded with chemically-bonded sand. EFFECTS OF Si CONTENT AND CE VALUE ON MICROSTRUCTURE The effect of Si content on nodule count for various CE values of 2 mm section green sand castings is shown in Fig. 7. As indicated in Fig. 7, for all three CE values, the nodule counts increase first, reach maxima at around 4.0-4.2%Si and then decrease slightly with a further increase in Si content. Also, the nodule count is higher for 4.15-4.35%CE (hypo- and near-eutectic) than for 4.45-4.55%CE and 4.75-4.85%CE (hyper-eutectic). On the other hand, Fig. 8 shows the changes in %ferrite with %Si for various CE values of 2 mm section green sand castings. As shown, the %ferrite increases continuously with increasing %Si in all three CE values, and also the irons with 4.15-4.35%CE exhibit somewhat higher %ferrite than the other two CE values.

Unit: mm

Fig. 4. Changes in nodule count and nodularity as a function of Si content for chemically-bonded sand castings containing about 3.0%C.

Fig. 5. Changes in %ferrite and %pearlite as a function of Si content for green sand castings containing about 3.0%C.


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(e)

(a) (b)

(c)

(f)

(g) (h)

(d)

Fig. 6. Microstructures of various heats, (a) A-HG2, (b) A-HG3, (c) C-HG2, (d) C-HG3, (e) G-HG2, (f )G-HG3, (g) J-HG2, and (h) J-HG3.


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In addition, the surface structures of some selected specimens (Heats D and G) were analyzed. The results show that no particular difference can be detected between the surface and the bulk in both 2 mm- and 3 mm sections cast in both green sand and chemically-bonded sand molds. There is an exception in one casting, G-HG3 (Heat G, High pouring temp, Green sand mold, 3 mm section), where a flake graphite rim of some 50-70 μm (0.05- 0.07 mm) can be observed. A series of photomicrographs depicting the surface structure of Heat G are shown in Fig. 9. The formation of the flake graphite skin can be attributed to the moisture in the mold and also to a relatively lower solidification cooling rate of that specific specimen compared to others.21-23 However, it has to be noted that in this study, the specimens used for thermal fatigue testing have been machined and also have a V-notch. Due to the fact that the initiation of cracks generally took place at the tip of the V-notch, the casting skin should not influence the thermal fatigue property in the present study. Nevertheless, in practical applications, the “skin-effect” does play an important role in affecting mechanical properties of parts which are used as cast without machining. This issue will be investigated in the authors’ future study.

EFFECT OF CHEMICAL COMPOSITION ON THERMAL FATIGUE LIFE Based upon the results concerning the influence of C and Si contents on microstructure (both nodule count and %ferrite), Heats A, C, G and J that contained about 3.0%C, but with various Si contents, were selected for thermal fatigue tests. The relationship between Si content (at a fixed C content of 3.0%) and thermal fatigue life (the total thermal cycles to failure) for castings with 2 mm and 3 mm sections and cast in green sand molds are respectively shown in Fig. 10a and b. As can be observed in Fig. 10, the thermal fatigue cycle to fracture increases first with increasing Si content, reaches maxima at around 4.6-4.8% Si, and then decreases substantially with a further increase in Si content. The trend of the change in thermal fatigue life (cycle) as a function of %Si is similar to that of %ferrite with %Si, but not that of nodule count with %Si. The results clearly indicate that the best thermal fatigue life corresponds to an iron microstructure that has the maximum percent ferrite rather than the highest nodule count. Accordingly, of the two primary factors that affect the thermal fatigue property, i.e., the nodule count and %ferrite, %ferrite plays a more important role than does nodule count in this respect. Based upon the obtained results, the optimal alloy design for attaining the best thermal fatigue property is a moderate hyper-eutectic composition, i.e., 4.5-4.6%CE, with a combination of relatively low C content and high Si content, e.g., 3.0%C + (4.6-4.8)%Si.

The effect of Mo on thermal fatigue property was studied by adding some 0.5%Mo to an iron with the alloy design corresponding to the highest thermal fatigue life as obtained above, Heat K. Data in Fig. 11 demonstrate the influence of Mo (0.5%) on thermal fatigue life of 2 mm section ductile cast irons cast in various conditions. Apparently, the number of thermal fatigue cycle to failure is substantially increased by introducing 0.5%Mo to the iron, regardless of casting conditions. Figure 12 illustrates the phenomenon of swelling near the thermal center of the thermal fatigue test specimen after failure. The change in specimen thickness, (in terms of increment in thickness) near the thermal center as a function of thermal fatigue cycle, is shown in Fig. 13. As seen in Fig. 13a, the increase in thickness is lower in an alloy with a higher Si content (4.77%) than with a lower Si content (2.98%). In addition, the alloy that contains 0.53%Mo exhibited less increase in thickness than that without Mo addition, as shown in Fig. 13b. An increase in thickness near the thermal center of the thermal fatigue test specimens is related to the stress relaxation at an elevated temperature. Generally, a more resistance to stress relaxation of an alloy would result in less increase in thickness. Therefore, alloys with a higher Si content and/or with Mo addition possess better resistance to stress.

Fig. 7. Effect of Si content on nodule count for various CE values of 2-mm green sand castings.

Fig. 8. Effect of Si content on %ferrite for various CE values of 2-mm green sand castings.


270

Fig. 9. Microstructures at casting skin of various specimens, (a) GHC2-100X, (b) GHC2-200X, (c) GHC3-100X, (d) GHC3-200X, (e) GHG2-100X , (f ) GHG2-200X, (g) GHG3-100X, and (h) GHG3-200X.

(e)

(a)

(c)

(g) (h)

(f)

(b)

(d)


271

Fig. 10. The effect of Si content on thermal fatigue cycle, together with %ferrite and nodule count, for castings containing about 3.0% C, and cast in green sand molds and at a higher pouring temperature (a) 2 mm section, and (b) 3 mm section. Fig. 11. The effect of Mo on total thermal fatigue cycle to failure.

Fig. 12. The shape of the thermal fatigue test specimens, (a) before test, and (b) after test.

(a)

(b)

(a)

(b)


272

This can explain the improved thermal fatigue property in alloys with a higher Si content and/or where Mo was added.

MICROSTRUCTURE TRANSFORMATION DURING CYCLIC THERMAL FATIGUE TEST (RT - 800C (1472F) For analyzing the microstructure transformation during the cyclic thermal fatigue tests within the temperature range of RT and 800C (1472F), we selected specimen G-HG2 (2 mm section cast at 1450C (2642F) into green sand molds) which contains 2.91%C and 4.77%Si (4.50%CE) and has ferrite content of 89%, nodule count of 881 #/mm2 and nodularity of 90%. The microstructure of specimen G-HG2 is shown in Fig.14a. For this specific alloy that contains 4.77%Si, the estimated critical temperature range is between 845 and 906C (1553-1663F). Within this critical temperature range, ferrite, austenite and graphite coexist. Due to the fact that the afore-mentioned critical temperature range exceeds the current maximum thermal cyclic temperature of 800C (1472F), theoretically it should be no phase transformation. However, because the maximum thermal cyclic temperature is quite close to the lower critical temperature, together with the phenomenon of Si segregation, phase transformation did take place in certain localized areas. Therefore, in the first thermal cycle, most of pearlite and part of ferrite in the as cast condition transformed to austenite during the heating and holding stage, and then the austenite transformed into martensite during the subsequent cooling cycle (water quench). As a result, martensite and un-transformed ferrite were present after the first thermal cycle, as shown in Fig. 14b. During the heating and holding stage, the dissociation of both cementites (in pearlite) and graphite nodules yields free carbon atoms, which diffuse to the vicinity of both the dissociated pearlite phase and graphite nodules. As a result, the carbon content in those areas is increased. The carbon-enriched areas favored the formation of austenite during the heating and holding stage, which in turn transformed into martensite upon cooling to room temperature. As a result, the transformed martensite is distributed mainly in areas immediately adjacent to graphite nodules and also in the vicinity of grain boundaries where the original pearlite phase was present. In the second thermal cycle, part of ferrite and martensite again transformed to austenite during the heating and holding stage, while at the same time the un-transformed martensite was tempered. In the subsequent cooling cycle, austenite once again transformed into martensite, and therefore, martensite, tempered martensite and un-transformed ferrite coexist in the microstructure. The afore-mentioned pattern of phase transformation continued to operate in the following thermal cycles. Fig. 14c and d, for example, show the microstructures of the specimens after 10 and 50 thermal cycles, respectively, where both specimens consist of ferrite, martensite and tempered martensite. However, at a

certain stage in the course of cyclic thermal fatigue test, minute graphite particles, the so-called secondary graphite, started to precipitate at grain boundaries, because of the repeated tempering of both martensite and tempered martensite. Figure 14e is an example of such precipitation of secondary graphite at grain boundaries of a specimen which had been run for 136 thermal cycles. The precipitation of secondary graphite has also been confirmed by Energy Dispersive X-ray Analysis (EDXA), as illustrated in Fig. 15.When the martensite was heated to the pre-set 800C (1472F), but failed to transform to austenite due to the temperature not being high enough for phase transformation to take place, the martensite was tempered at a relatively high temperature 800C (1472F) instead, causing the precipitation of secondary graphite along the grain boundaries. With the progress of the thermal cycles, the precipitation of the secondary graphite also continued, which resulted in a gradual reduction in the dissolved carbon content in the matrix. Consequently, the volume fraction of the ferrite phase increases, while at the same time the amounts of martensite and tempered martensite decrease. Figure 14f is the microstructure of the specimen that had been tested for 766 cycles and finally fractured, which consists of mainly ferrite and numerous graphite particles precipitated at the grain boundaries. Under the influence of residual tensile stress developed in each thermal cycle, cracks, once they occur, were initiated at the graphite-matrix interface and/or in the vicinity of the precipitated secondary graphite particles. Figure 16 shows the Scanning Electron Microscope (SEM) micrograph for the case of 136 thermal cycles of specimen G-HG2, while Fig. 17 shows the microstructure of the finally failed specimen, where secondary graphite particles and cracks along the grain boundaries are quite evident. CONCLUSIONS In this study, the correlations between alloy composition, microstructure and thermal fatigue life were analyzed for attaining the optimal alloy design and casting conditions. Based upon the results obtained herein, the following conclusions can be drawn. 1. For a fixed C content of some 3.0%, both the nodule

count and graphite nodularity increase first with increasing Si content, reach maxima at around 3.8-3.9%Si and then decrease with further increase in Si content. On the other hand, the percent ferrite increases gradually with increasing Si content, reaches a maximum at around 4.6-4.8%Si, and then remains more or less constant or decreases slightly. In addition, irons with 4.45-4.55%CE exhibit somewhat higher nodule count and %ferrite than those with 4.75-4.85%CE.


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Fig. 13. (a) A comparison in the increment of thickness as a function of thermal cycle between alloys with different Si contents. (b) A comparison in the increment of thickness as a function of thermal cycle between alloys with and without Mo addition. Fig. 14. Microstructure changes during thermal fatigue test of specimen G-HG2, (a) as-cast, (b) 1 cycle, (c) 10 cycles, (d) 50 cycles, (e) 136 cycles, and (f) 766 cycles and fractured.

(a)

(f)

(b) martensite

ferrite

(c)

temper martensite

martensite (d)

martensite

temper martensite

(e)

Secondary graphite


274

Fig. 15. The EDXA analysis of the secondary graphite particle. 1 Fig. 16. The SEM microstructure of specimen G-HG2 after 136 thermal cycles.

Fig. 17. The SEM micrographs of specimen G-HG2 after fracture.

EDXA

secondary graphite

cracks


275

2. Regarding the effects of casting parameters on microstructure, the results show that higher nodule counts can be obtained in castings with a thinner section or with a lower pouring temperature. However, the effect of molding material (chemically-bonded sand and green sand) on nodule count is not significant. Furthermore, the percent ferrite is higher in castings with a thicker section, a higher pouring temperature and molded with chemically-bonded sand.

3. The optimal alloy design for attaining the best

thermal fatigue property has been found to be a moderate hyper-eutectic composition, i.e., 4.5-4.6%CE, with a combination of relatively low C content and high Si content, e.g., 3.0%C + (4.6-4.8)%Si. In addition, the best thermal fatigue property corresponds to a microstructure having the highest ferrite content and moderate nodule count.

4. Adding some 0.5%Mo to the iron significantly

increases the thermal fatigue life. Also, the phenomena of microstructure changes and the failure mechanism during the cyclic thermal fatigue process were also analyzed.

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