Microalloying in Austempered Ductile Iron (ADI)

Transcription

1 Paper pdf, Page 1 of 12 Copyright 2012 American Foundry Society Microalloying in Austempered Ductile Iron (ADI) D.S. Padan Tata Motors Limited, Jamshedpur, India ABSTRACT Austempered Ductile Iron (ADI) is well known for its outstanding engineering properties such as tensile strength, elongation, wear resistance etc. which are comparable with those of steel. Normally ADI is alloyed with elements like Mo, Ni, Cu and Mn. Microalloyed steel is another successful development, where presence of elements like V (vanadium) and Nb (niobium) significantly improve its mechanical properties. Based on the experience gained in microalloyed steel, it was expected that beneficial effects of microalloying elements, V and Nb, may also be exploited in ADI as well. Accordingly, experiments were carried out on microalloyed ADI. The effect on mechanical properties, microstructure, austemperability and wear characteristics was examined. The study confirmed that microalloying of ADI with 0.1 % V or % Nb with 1.4 % Ni improved ultimate tensile strength significantly and it was superior to that of conventional ADI alloyed with 0.3 % Mo and 1.4 % Ni. The wear resistance of microalloyed ADI improved remarkably. No significant change in austemperability was observed. Improvements in mechanical properties were attributed to presence of precipitates of V and Nb carbides. Keywords: austempered ductile iron, microalloyed INTRODUCTION Austempered ductile iron (ADI) is an emerging engineering material with wide application in the automotive and other industries. ADI possesses an excellent combination of mechanical properties like high tensile strength ( MPa), elongation (3-15 %), toughness (60-90 MN -3/2 ), wear resistance and fatigue strength. 1,2,3,4,5 A number of forged steel components have now been replaced by ADI castings, especially in automotive applications. These include crankshafts, camshafts, differential gears, chain links, spiral, ring and pinion gears and universal joints etc. 6 ADI is normally alloyed with Mo, Ni, Mn and Cu in suitable combinations. 7 These elements improve hardenability and assist in stabilizing austenite. In view of the attractive effects of V and Nb on transformation in steels, it was speculated that microalloying additions of these elements to ductile iron may also be beneficial. Previous studies on the effects of microalloying with niobium and vanadium have been conducted only on ferrite - pearlitic grades of ductile irons. 8,9 It was anticipated that the benefits of additions of V and Nb would justify microalloying ADI with these elements. Therefore, the objective of the present investigation was to study the effects of microstructure, mechanical properties, wear characteristics and austempering response when microalloying ADI with V and Nb and to compare the results with that of standard ADI containing Mo and Ni. EXPERIMENTAL PROCEDURE MELTING The liquid iron was produced in the coreless induction furnace in a commercial foundry from steel scrap, foundry returns, pig iron, petroleum coke and ferro-silicon charges. The chemical composition of the liquid iron was adjusted for typical grey cast iron containing % C, % Si, 0.35 % max. Mn, 0.06 % max. S, 0.04 % max. P and 0.07 % max. Cr etc. The melt was tapped at a temperature of 1490C (2714F). Magnesium Treatment Carburizer (graphite) was added to liquid metal from the furnace and the melt was treated with pure magnesium lumps in a Georg Fischer converter (1.5 Ton capacity). The quantity of magnesium added was estimated per following expression. 10 Mg = 3/4(Initial S content) + * Residual Mg ( %) ** Expected Mg recovery *Aimed-0.04% **50% Ladle Addition and Casting The magnesium treated melt was tapped into the experimental ladle (100 Kg capacity) along with post inoculants Fe-Si and misch metal (cerium). Estimated quantities of Ni (Pure), Fe-V (60.0 % V), Fe-Nb (50.0 % Nb) and Fe-Mo (45.0 % Mo) were added along with liquid metal stream in the ladle to produce the experimental alloys with the specified chemical compositions. Metal stream agitation during ladle filling ensured proper mixing of the alloying elements that were added. The melt was cast into standard 30 mm (1.18 in.) test bars in oil sand molds. AUSTEMPERING TREATMENT The test bars, packed in charcoal (to avoid oxidation), were austenitised at C ( F) for 1 hr in a pit furnace, followed by transferring the bars rapidly into

2 Paper pdf, Page 2 of 12 the salt bath (NaNO 3 -KNO 3 ) maintained at 335C (635F), for austempering. Initially, to study the progress of the austempering reaction, small test pieces (18 Φ x 18 mm) machined from the test bars were subjected to austempering treatment for different times ranging from about 5 min to 3 hr. These few alloys that developed enhanced mechanical properties after austempering were later subjected to prolonged austempering at the same temperature for different times (beyond 3 hrs) up to 16 hrs. The hardness values of the austempered test pieces were noted. The optimum austempering time was selected for which the austempered specimens developed the minimum hardness values with maximum ausferrite. The same austempering schedule was carried out on the test bars for making tensile and wear specimens. The bars were austenitized at 900C (1652F) for 1 hr, followed by austempering at 335C (635F) for 1.5 hrs. TENSILE TEST SPECIMEN Tensile specimens were fabricated from the as cast test bars as per standard IS to measure the tensile strengths of the alloys in as cast condition. The austempered test bars were used to fabricate tensile specimens as per standard JIS No. 4 to measure the tensile strengths of the austempered ductile irons. WEAR TEST SPECIMEN Wear specimens of dimension 10 mm (0.39 in.) diameter x 30 mm (1.18 in.) were fabricated from the austempered test bars. TENSILE AND HARDNESS TESTS A minimum of two specimens were tested for tensile strength for each sample. The hardness values in BHN were measured with 5 mm (0.20 in.) indenter under 750 kg load and the average of three readings was noted. X-RAY DIFFRACTION STUDY To estimate the austenite content, the austempered ductile irons samples were studied with the Philips PW 1820 X- ray diffractometer using CoKα radiation. The equipment was operated at 40KV and the samples were scanned through 45 to 113, 2 θ values. The areas under the intensity peaks of diffractions from austenite and ferrite planes were measured and the volume fractions of austenite were estimated. X-Ray Diffraction (XRD) technique was also used to detect the alloy carbide precipitates in microalloyed ADIs, but due to their very low volume fractions, the corresponding intensity peaks were not found. WEAR TEST Adhesive wear tests were carried out with a pin on a disc type wear testing machine. The wear specimen was held vertically against the rotating disc (HRC 65) under different loads for one hour. The disc speed was maintained at 215 rpm and the track diameter was 176 mm (6.93 in.). The corresponding linear wear speed for the specimen was 1.98 m/sec. The tests were carried out under loads of 29.4 N, 39.2 N and 49.0 N. The height loss of the specimen due to wear in micrometer (μm) and frictional force in Kgf were noted from the digital monitor fitted with the equipment. From this data linear wear rate, coefficient of friction and average coefficient of wear resistance were calculated. MICROHARDNESS MEASUREMENT Microhardness (HV0.025) values of the unworn and the worn out surfaces were measured under 25 gm load for 15 sec loading time and the average of five readings were noted. MICROSCOPY Unetched samples of as cast as well as austempered ductile irons were examined under optical microscope to study the microstructural characteristics, like nodularity and nodule count. To study the matrix microstructures of as cast and austempered ductile irons, the polished samples were etched with 2% Nital. Some selected ADI samples were studied under Scanning Electron Microscopy (SEM) (JEOL, Model-JSM-5800) and the qualitative analysis of the distribution of alloying elements of the iron were done by Energy Dispersive X-ray analysis (EDX). The worn surfaces and the corresponding debris of the selected ADIs were examined under SEM. RESULTS EXPERIMENTAL ALLOYS Ductile iron alloys used in the present investigation have been detailed with their designations and final chemical composition in Table 1. B is base alloy containing 1.4 % Ni only. V has 0.1 % V and 1.4 % Ni, alloy N has % Nb and 1.4 % Ni and alloy M has 0.3 % Mo and 1.4 % Ni. M has been taken as the standard alloy for comparison in the present study. As the alloys V, N and M have 1.4% Ni in common, these are simply designated mostly as alloy V (0.1% V), alloy N (0.043% Nb) and alloy M (0.3% Mo), respectively, in the following text. As alloy U is a plain ductile iron without any alloying element, it is not discussed much in the present study The nodularity, nodule count, ferrite percentage, tensile strength, yield strength, percent elongation and hardness values of alloys in as cast condition have been indicated in Table 2. From the data, it is observed that the percentages of nodularity of all the alloys are in between 85 to 90 %. The corresponding nodule counts /mm 2 vary in between 200 to 304. The percentages of ferrite in the matrices of ductile irons alloyed with V, Nb, Mo along with Ni are less than that

3 Hardness (BHN) Paper pdf, Page 3 of 12 Table 1. Chemical Compositions of Ductile Irons Investigated Chemical composition % C % Si % Mn % Ni % V % Nb %Mo U B V N M Table 2. Mechanical Properties of As Cast Ductile Irons Investigated Nodularity Nodule count Matrix Ferrite / mm 2 (%) Tensile strength (MPa) Yield strength (MPa) Elongation (%) Hardness (BHN) U B V N M present in the unalloyed ductile iron (alloy U) or in the iron alloyed with Ni only (alloy B). This is also supported by their high hardness values as compared to those of alloys U and B. The tensile strengths of alloys V and N are 19 to 22 % more than that of alloy B. The increase in the pearlite content and the corresponding increase in the hardness values and tensile strengths of the as cast ductile irons alloy V and N can be attributed to the formation of the complex eutectic carbides or simple alloy carbides of V and Nb during solidification AUSTEMPERING HEAT TREATMENT To study the austempering behavior of the ductile iron alloys under investigation, the alloys were initially austenitized at 900C (1652F) for 1 hr and austempered at 335C (635F) for different austempering times ranging from 0.5 to180 min followed by air cooling. The trend of variation in hardness (BHN) is indicated in Fig. 2. The results obtained indicate that microalloyed ADIs (alloys V and N), the standard ADI (alloys M) and base ADI (alloy B) follow the same pattern of hardness variation. It indicates that as the austempering time increases, the corresponding hardness value decreases drastically up to 60 min of austempering time and the same becomes almost constant up to 180 min (except for alloy B). The minimum hardness values are obtained after austempering for 90 min. In case of alloy B, the hardness value increases sharply after austempering for 150 min. The hardness value of alloy U does not follow the usual trend of austempering reaction as observed with other alloys. Since ADIs alloy V (0.1 % V) and alloy N (0.043 % Nb) developed superior mechanical properties (discussed later) as compared to those of the standard ADI (alloy M) containing 0.3 % Mo, alloys V, N and M were subjected to prolonged austempering up to about 960 min at the same austempering temperature in order to compare their austempering kinetics. The trend of hardness developed after prolonged austempering is shown in Fig. 3. These figures indicate that the hardness values, developed after austempering for 60 min, remain almost constant up to 240 min of austempering; followed by a sharp increase in the hardness and continued to increase up to 960 min of austempering time. Figure 4 presents the XRD pattern of alloy V austempered for 90 and 960 min Austempering time (Minute) U B V N M Fig. 2. This graph shows the variation in hardness of ADI alloys U, B, V, N and M versus austempering time up to 180 min.

4 UTS (MPa), Hardness (BHN) Elongation (%) Hardness (BHN) Paper pdf, Page 4 of Table 3. Mechanical Properties of Austempered Ductile Irons (Austenitized at 900C [1652F] for 1 Hr; Austempered at 335C [635F] for 1.5 Hr) Austempering time (Minute) V N M Tensile strength (MPa) Yield strength (MPa) Elongation (%) Hardness (BHN) U B V N M Fig. 3. This graph shows the variation in hardness of ADI alloys V, N and M versus austempering time up to 960 min UTS UTS BHN E 200 BHN 2 0 U B V N M 0 Fig. 5. This graph shows the ultimate tensile strength (UTS), hardness and elongation of alloys U, B, V, N and M. Fig. 4. The graph shows the XRD pattern of alloy V austempered at 335C for (a) 90 min and (b) 960 min. MECHANICAL PROPERTIES The mechanical properties of austempered ductile irons are listed in Table 3 and depicted in graphical form in Fig. 5. Representative optical photomicrographs of the ADIs are shown in Figs. 6a-6d. Representative SEM photomicrographs associated with element distribution mappings of alloys V, N and M are shown in Figs. 7 to 9, respectively. The tensile properties (tensile strength, elongation) of representative ADIs investigated with respect to those of six grades of ADI specified in ASTM A897/ A897M-06 are indicated in Fig The percentage increase in ultimate tensile strength (UTS) of ADIs alloy V (0.1 % V), alloy N (0.043% Nb) and alloy M (0.3 % Mo) with respect to that of ADI alloy B are 17 %, 5 % and 3 %, respectively. The percentage decrease in elongations of ADI alloys V, N and M with respect to ADI alloy B are 19 %, 25 % and 48 %, respectively. The tensile strengths of ADI alloys V and N are 13 % and 2 % higher than that of the standard ADI alloy M, respectively. The tensile properties of alloys V and N are close to grade III and II respectively with marginally low hardness in comparison to ASTM standard. (c) X 480 (d) x 480 Fig. 6. (a-d) Optical photomicrographs of austempered ductile irons show: (a) base alloy B (1.4% Ni only), (b) alloy V (0.1% V), (c) alloy N (0.043% Nb) and (d) standard alloy M (0.3% Mo.)

5 Tensile strength (MPa) Paper pdf, Page 5 of 12 Fig. 7. Element distribution mappings for V and Ni in ADI alloy V (0.1% V) show: (a) elected area for mapping (500x), (b) Vanadium distribution and (c) Nickel distribution Fig. 9. Element distribution mappings for Mo and Ni in ADI alloy M (0.3% Mo) show: (a) selected area for mapping (500x), (b) Molybdenum distribution and (c) Nickel distribution ASTM A897/A897M ALLOY M ALLOY V ALLOY N Elongation ( % ) Fig. 10. This graph shows the comparison of the tensile properties of ADI alloys V, N and M with the six ADI grades specified in standard ASTM A897/ A897M Fig. 8. Element distribution mappings for Nb and Ni in ADI alloy N (0.043% Nb) show: (a) selected area for mapping (500x), (b) Niobium distribution and (c) Nickel distribution

6 Microhardness (VPN) Cumulative wear loss (micro-m) Ave. linear wear rate x 1/1000 (micro-m/m ) Linear wear rate x 1/1000 (micro- Paper pdf, Page 6 of 12 WEAR CHARACTERISTICS Cumulative Wear Loss and Linear Wear Rate Figure 11 represents a sample in graphical form of cumulative wear loss in µm during the wear test of ADI alloy V under a wear load of 39.2 N. The linear wear rate in micrometer per meter (µm/m) is calculated as the height lost due to wear divided by the corresponding sliding distance ( h/ l) at the interval of 3 min from the start of wear tests. It is observed that wear of specimens occur in two stages. In the initial stage, the wear rate is high, which gradually lowers and becomes steady. One sample graph representing the linear wear rate trend for ADI alloy V under 49.0 N load is shown in Fig. 12. Table 4 indicates the average linear wear rates in µm/m, calculated for the distance of about 3.6 km during the steady state of wear. The same is represented graphically in Fig. 13. The average linear wear rate (ALWR) of alloys V (0.1% V) and N (0.043% Nb) are lower by 25 % and 8 %,respectively, compared to that of alloy B under 29.4 N wear load which is further lowered by 34% and 26 %, respectively, under 49.0 N wear load. Similarly, the ALWR of alloys V and N are lower than that of alloy M (0.3 %Mo) which are 16 % and 5 % under a 49.0 N wear load. The ALWRs of all alloys increases with an increase in wear load. The variation of microhardness of the worn out surface of ADI under different wear loads is shown in Table 5. The same is represented graphically in Fig. 14. The microhardness of the worn out surface of the ADI specimen is high with respect to that of the unworn surface and the same increases with the increase in wear load Fig. 11. This graph shows the cumulative wear loss versus sliding distance under 39.2 N wear load for ADI alloys, B, V, N and M. Table 4. Average Linear Wear Rates (ALWR) of ADIs Investigated Under Different Wear Loads Sliding distance x 119 (m) ALWR x 10-3 in µm/ m 29.4 N load 39.2 N load 49.0 N load B V N M B V N M 0m/m) Sliding distance x 119 (m) Fig. 12. This graph shows the variation in linear wear rate with sliding distance under 49.0 wear load for ADI alloy V B V N M 29.4 N Load 39.2 N Load 49.0 N Load Fig. 13. This graph shows the average linear wear rate of ADI alloys B, V, N and M under different wear loads. Table 5. Microhardness Values of Unworn and Worn Out Surface of ADIs Under Different Wear Loads Microhardness (HV) Unworn Worn out surface surface 29.4 N load 39.2 N load 49.0 N load B V N M B V N M before w ear 29.4 N Load 39.2 N Load 49.0 N Load Fig. 14. This graph shows the variation in microhardness of unworn and worn surfaces of ADI alloys B, V, N and M under different wear loads.

7 Ave. coeff. of wear resistance (N m /micro-m) Paper pdf, Page 7 of 12 Average Coefficient of Wear Resistance The average coefficient of wear resistance (ACWR) expressed in Newton (N)-meter (m) / micrometer (µm) is defined as average frictional force times distance traveled divided by height lost due to wear. 12 The calculated values of ACWR of ADIs travelling a distance of 3.6 km in the steady state under different wear loads are indicated in Table 6 and in graph (Fig 15). ADI alloy V (0.1% V) has the maximum value of ACWR followed by ADI alloys N, M and B. A Comparison of ACWR values of ADIs under 49.0 N wear load indicates that alloy V has wear resistance of about 1.56 and 1.15 times those of alloys B and M, respectively. N has wear resistance of about 1.35 times that of alloy B and almost the same as that of alloy M. M has wear resistance about 1.35 times that of alloy B. The ACWR values of ADIs that are examined have the tendency to decrease when the applied wear load is increased. corresponding SEM photographs for alloy M are shown in Figs. 18 a, 18b and 19, respectively. The SEM photographs of the worn out surface and corresponding wear debris of alloy V exhibit mostly the growth of oxides on the worn out surface and the oxide debris, respectively. The corresponding photographs for alloy M show the presence of cracks and delamination of the worn out surface along with oxide layers formed at some localized regions and the metallic chips as wear debris, respectively. Table 6. Average Coefficient of Wear Resistance (ACWR) of ADIs under Different Wear Loads ACWR in Nm/ µm 29.4 N load 39.2 N load 49.0 N load B V N M Fig. 16. This is a SEM photograph of worn surface of ADI alloy V (0.1% V), subjected to wear under 49.0 N load N Load 39.2 N Load 49.0 N Load B V N M Fig. 15. This graph shows the average coefficient of wear resistance (ACWR) of ADI alloys, B, V, N and M, under different wear loads. Mechanism of Wear It is observed that microalloyed ADI alloy V (0.1 % V) is more wear resistance than the other alloys investigated. Therefore, this alloy has been examined in details for studying its wear mechanism through SEM and the same has been compared with that of the standard ADI alloy M (0.3% Mo). Figures 16a, 16b and 17 show the SEM photographs of the surface of alloy V worn under 49.0 N load and the corresponding wear debris, respectively. The Fig. 17. This is a SEM photograph of wear debris of ADI alloy V (0.1% V), generated during wear test under 49.0 N load.

8 Paper pdf, Page 8 of 12 Fig. 18. This is a SEM photograph of a worn surface of standard ADI alloy M (0.3% Mo), subjected to wear under 49.0 N load. Fig. 19. This is a SEM photograph of wear debris of standard ADI alloy M (0.3% Mo), generated during wear test under 49.0 N load. DISCUSSION AUSTEMPERING HEAT TREATMENT The observed pattern of variation of the hardness values of ADIs with austempering time can be explained on the basis of the published information, regarding the progress of austempering reaction that occurs through different stages. The austempering reaction, after the spinodal decomposition of initial austenite into low carbon and high carbon austenites, follows the stage I reaction in which the low carbon austenite undergoes bainitic transformations. 13 The initial high hardness value is due to the fact that with short austempering time less amount of the low carbon austenite has been transformed into ausferrite (bainitic ferrite + high carbon austenite) and therefore, the majority of the same has been transformed into martensite on subsequent air cooling. As the austempering reaction progresses, the low carbon austenite gradually transforms to ausferrite and the amount of untransformed low carbon austenite converting into martensite on air cooling reduces and therefore, the corresponding hardness values fall drastically. The low hardness values including the minimum one correspond to stage II reaction, in which the reaction products exhibit an ausferrite structure with negligible or no martensite. Subsequently, the increased hardness values after 360 min of austempering for alloys V, N and M and 150 min for alloy B are due to the start of stage III reaction, when the austenite starts decomposing into ferrite and carbide on prolonged aging. Dubensky et al. also reported the same trend of hardness variation with progress of austempering reactions. 14 Figure 4 exhibits representative XRD pattern for alloy V indicating the reduced intensity peaks from austenite (2,0,0) and (2,2,0) planes after austempering for 960 min with respect to that developed after austempering for 90 min, which is due to decomposition of austenite into ferrite and carbide during prolonged austempering. Based on the variation of hardness with the austempering time, an attempt has been made to develop the graphical representation, exhibiting the austempering reaction and processing windows for the ductile irons microalloyed with 0.1 % V (alloy V), % Nb (alloy N) and the standard iron alloyed with 0.3 % Mo (alloy M) in order to compare their austempering behaviors. It is assumed that: (a) minimum bainitic transformation has taken place at the minimum austempering time when the hardness values are maximum; (b) maximum bainitic transformation has taken place by 60 min of austempering time when corresponding hardness values are minimum and remain almost the same with further increase in austempering time; (c) minimum decomposition of austenite has taken place after austempering for 240 min beyond this time only the hardness values increase sharply and (d) major quantity of austenite decomposition has taken place by 960 min of austempering time. The percentages of bainitic transformation between 0.5 to 60 min and the percentages of decomposition of austenite between 240 to 960 min of austempering time are the percentage of their corresponding maximum values. Figures 20a and 20b show the progress of the austempering reaction for alloys V, N and M with normal austempering time and log (austempering time), respectively, based on the assumptions made. A comparison of the austempering reactions from this figure indicates that the difference in the safe time interval (t 2 -t 1 ) or processing window for the microalloyed ductile iron alloys V and N with respect to that for standard iron alloy M are not much different. A close observation rather indicates that the alloy V responds better to bainitic transformation as compared to alloy M. Therefore from the above comparison, it seems that the iron microalloyed with 0.1 % V along with Ni (alloy V) and % Nb along Ni (alloy N) can be austempered successfully under conditions adopted for standard iron alloyed with 0.3 % Mo along with Ni (alloy M) to develop the optimum mechanical properties. The unusual hardness variation pattern, observed with alloy U, is most likely due to its inadequate hardenability in the absence of any

9 Paper pdf, Page 9 of 12 hardenability promoting element. This makes the alloy mostly unsuitable for austempering, except for thin sections. ALLOY V N M MECHANICAL PROPERTIES A review of published literature indicates that additions of Ni up to 1 to 1.5 % have a slight effect on the tensile strength and hardness of the ductile irons austempered at temperatures between 300 to 400C ( F). 10 Nickel has tendency to suppress carbide formation in lower ausferrite, resulting in an increased ductility. The same explanation can be put forward for explaining practically the same tensile strength and hardness values and increased elongation of the ADI alloyed with Ni only (alloy B) as compared to those of the unalloyed ADI (alloy U). It is observed that microalloyed ADI alloys V (0.1% V) and N (0.043% Nb) and the standard ADI alloy M (0.3% Mo) possess increased strengths with slightly reduced ductility as compared to those of ADI alloy B with Ni only. These improvements in the mechanical properties can be attributed to one or more of the following strengthening mechanisms: (a) solid solution strengthening, (b) grain refinement and (c) precipitation hardening through the precipitation of fine carbides of the alloying elements added. ALLOY V N M AUSTEMPERING TIME (Minutes) (a) LOG AUSTEMPERING TIME (Minutes) Fig. 20. Graphs show progress of austempering reactions at 335C for alloys V, N and M with (a) austempering time and b) log austempering time. (b) The alloying elements like V and Nb are widely used for developing the high strength low alloy (HSLA) steels, because they have potentials for precipitation strengthening and pronounced grain refining even when added in small quantities. The enhanced mechanical properties of the HSLA steels are obtained partly by the reduced ferrite grain size due to pinning of ferrite grain boundaries by the alloy carbide or carbonitride precipitates and partly by the precipitation strengthening effect. For ductile irons, V has been considered an undesirable element due to its tendency to promote the formation of eutectic carbides or chill. However, Dawson reported that the undesirable eutectic carbides containing V can be removed by the heat treatment of ductile irons, and, therefore, improve their strengths and ductility through the formation of the fine vanadium carbide or carbonitride precipitates by the mechanisms similar to those observed in HSLA steels. 15 Shen et al. studied annealed ductile irons alloyed with small quantity of vanadium. 16 They noted that the combinations of intermediate strengths with good elongations of annealed ductile irons were developed, which were not due to the changes in the graphite morphology or volume fraction of nodules. They were due to the reduced ferrite grain size occurring through the pinning effect of grain boundaries due to the presence of fine precipitates of vanadium carbide (V 4 C 3 ), formed through dissolution and re-precipitation of vanadium containing coarse eutectic carbides during heat treatment. The same mechanism can be put forward here to explain the improved mechanical properties of ADI alloy V (0.1% V) with respect to those of ADI alloy B with Ni only. The only difference can be that quantity of the precipitate formed can be less, as the time for

10 Paper pdf, Page 10 of 12 precipitation is less with the austempering process than with the annealing process. The element distribution mappings for V and Ni present in microalloyed ADI alloy V (0.1% V), as seen in Fig. 7, reveal that both the elements are almost uniformly distributed throughout the matrix of the microstructures except the graphite nodules. A very close observation of the vanadium distribution mapping reveals that its distribution is mostly circular in shape, which is most likely due to the precipitation of vanadium carbide (V 4 C 3 ) along the grain boundaries of the austenite- ausferrite, formed during the austempering heat treatment as observed with heat treated HSLA steels and with annealed ductile irons containing vanadium. The presence of vanadium carbide precipitates probably resists the bainitic ferrite grain growth by pinning the grain boundaries. This results in grain refinement and thereby increases the strength of ADI microalloyed with vanadium. The other possible reasons for the enhanced mechanical properties are: precipitation hardening by vanadium carbide precipitates and the presence of undissolved fine vanadium (eutectic) carbides in the matrix. The effects of niobium in high strength low alloy steels or in cast irons are reported to be similar to that of vanadium as both of them have the tendencies for the formation of carbides / carbonitrides during solidification and cooling. 17,18 The only difference is in the amount of these compounds, formed due to their differences in kinetics, the temperature of precipitation reaction and the extent of dissolution of carbides during reheating for heat treatment. The dissolution of niobium carbides / carbonitrides takes place at very high temperatures (about 1150C [2102F]) as compared to that of coarse vanadium carbides (about 1050C [1922]F). Therefore, the degree of re-precipitation of niobium carbide precipitates may be much less as compared to that of the vanadium carbide precipitates and, therefore, it contributes less for strengthening through precipitation in ADI. The presence of small amounts of niobium also retards the austenite grain growth. 17 The element distribution mapping for Nb and Ni present in microalloyed ADI alloy N (0.043% Nb) as shown in Fig. 8 reveal that both niobium and nickel are distributed in similar fashions as in case of vanadium and nickel in ADI alloy V. The distribution of niobium in the matrix is very scattered which is probably because of the low percentage of niobium (0.043%) present in the alloy. The analysis of the above observations indicates that the high strength of ADI containing niobium is most likely due to the presence of undissolved niobium carbides / carbonitrides and partly due to the reduced grain size as well as the precipitation hardening effects due to the formation of fine particles of niobium carbide precipitate. Molybdenum is reported to increase and adjust the hardenability of the irons and steels. 10 It segregates strongly during solidification and forms carbides. The tendency of segregation becomes more pronounced if its percentage exceeds 0.3 %. The published information indicates that the tensile strength, hardness and percentage elongation of ADI decrease progressively as the Mo content is increased. 19 Therefore, the low improvement in tensile strength of the standard ADI alloy M (0.3% Mo) over ADI alloy B, containing Ni only, is probably due to their combined effects of controlling the austemperability and the resulting efficient bainitic transformation of the iron and the presence of undissolved carbides, formed during solidification. Figure 9 shows the element distribution mapping for molybdenum and nickel in the standard ADI alloy M (0.3% Mo). This pattern of distribution is most likely due to the segregation of eutectic carbide containing Mo along the eutectic cell boundaries, formed during solidification. These facts of grain refinement and precipitation hardening explain the development of improved mechanical properties of microalloyed ADIs with even 0.1 % V (alloy V) and % Nb (alloy N) with respect to that of the standard ADI alloyed with 0.3 % Mo (alloy M). WEAR CHARACTERISTICS The initial wear rates are high due to the initial rough surfaces of the specimen in the unworn condition developed during machining. This will develop high frictional force as well as high stress concentrations on the asperities of the surface under the applied wear load. This results in microwelding and their ruptures and therefore leads to the removal of the protruded surfaces in the form of debris. The subsequent lower and steady wear rate is due to smoothening of the specimen s surface and the strain hardening effect of austenite in the ausferrite structure. 20 This phenomenon is also supported by the increase in microhardness values of the worn out surfaces with respect to those of the corresponding unworn surfaces as shown in Table 5 and Fig. 14. This stage is known as steady state of wear. Another additional cause for lower wear rate can be the lubrication of the wearing surface due to smearing effect of the exposed graphite nodules or the formation of the loose iron oxides layer on the surface. The evidence of fine graphite layer with a lustrous grayish color was observed on the rotating disc surface of the equipment after the test. It is reported that the presence of even small quantity of fine and hard particles of vanadium carbide and niobium carbide or carbonitrides in the iron increase its wear resistance. 21 Therefore, it can be explained that the lower wear rates of ADI alloys V (0.1% V) and N (0.043% Nb) with respect to that of ADI alloy B with Ni only are mostly due to the presence of hard and fine precipitates of vanadium carbide and niobium carbide, respectively, formed during the austempering heat treatment and also due to the presence of undissolved fine eutectic carbides containing vanadium and niobium as previously discussed. The presence of these hard and uniformly

11 Paper pdf, Page 11 of 12 distributed fine alloy carbides precipitates, embedded in tough wearing surface, act as barrier against wear until they are pulled out due to the continuous sliding under load. Molybdenum is known to enhance the austempering kinetics when added with nickel. It has the tendency to segregate in the intercellular region and form carbides during solidification. 10 Therefore, the standard ADI alloy M (0.3% Mo) has resistance against wear because of its better austemperability and also due to some eutectic carbide formed during solidification and therefore, has a lower wear rate than that of alloy B but more than those of microalloyed ADI alloys V and N. The increase in average linear wear rates of ADIs examined under the increasing wear load is due to the combined effects of: (a) increased frictional force, leading to easy nucleation and propagation of the surface cracks and subsequently removing the surface layers and (b) increase in frictional temperature, resulting from the formation of a greater amount of wear debris of iron oxides. The experimental data indicate that the increase in the wear rates of the microalloyed ADI alloys V (0.1% V) and N (0.043% Nb) under the increasing wear loads is less as compared to those of the ADI alloy B and the standard ADI alloy M. It means that the presence of carbides of V and Nb in the matrices increases the load bearing ability of the corresponding alloys and thereby improves their wear resistances even under the increased wear loads. The calculation of average coefficient of wear resistance (ACWR) involves the height loss due to wear; the frictional force which depends on the coefficient of friction for a particular applied wear load; and the sliding distance traveled. Therefore, the variation of ACWR values as observed are due to the variation of these parameters. The high ACWR values of microalloyed ADI alloys V (0.1% V) and N (0.043% Nb) and of the standard ADI alloy M (0.3%Mo), as compared to that of ADI alloy B containing Ni only, are basically due to lower wear rates and lesser height loss of the former alloys as compared to that of the later for the same distance traveled under a particular applied wear load. The decrease in ACWR value of particular ADI with the increase in the wear load is mainly due to the corresponding increase in height loss because of high stress and high frictional temperature developed at the surface. A review of published literature indicates that the wear of ADIs can occur through one or more of the following mechanisms: (a) oxidative wear, (b) stick and slip (microweldings and rupturing) wear and (c) lamination and delamination wear. Sastry, in his research work on the mechanism of wear of ADIs, has indicated that oxidative wear mostly occurs during adhesive type of wear of ADI. 22 It is also reported that the iron oxides formed on the cast iron are not tenacious and can be removed easily during sliding. During wear testing, the temperature of the contact surface rises considerably (150C [302F]) and the surface layer becomes defective due to the mechanical deformation. Therefore, the diffusion rate of oxygen through the defective surface layer can be quite high which results in the formation of oxides and consequently leads to oxidative wear. The wear of ADI can also occur due to fatigue phenomenon. 23 This consists of the nucleation of cracks below the deformed surface, followed by their growth during wearing that finally leads to the breakage of surface layers. In steels, the nucleation of cracks is the controlling factor for wear and the delamination theory is more generally accepted for wear of steels. In ADI, crack nucleation is easier than in steels owing to the presence of graphite nodules. Graphite nodules close to the wearing surface deform readily and become ellipsoid in shape and therefore, act as cracks just below the wearing surface of the specimen, which subsequently grows. 5 The cracks grow parallel to the surface under loading until they reach a length, which results in a complex state of stress that causes the crack tip to change its path toward the surface and then the final breakage of the metal layers immediately follows. On the basis of the published information as previously summarized and the examination of the SEM photographs of worn surfaces developed under 49.0 N wear load and the corresponding wear debris, it is evident that ADI alloys V and M wear out through mixed mechanism of oxidative and delamination types of wear. The observations indicate that the presence of fine and hard precipitates of alloy carbides in microalloyed ADIs do not alter the basic mechanism of wear of ADI. Therefore the increased wear resistance of microalloyed ADIs investigated is mostly due to increase in their load bearing abilities due to the presence of alloy carbides in the tough ausferrite matrices. CONCLUSION A study on austempering response, mechanical properties and wear characteristics of ADIs microalloyed with 0.1 % V or 0.043% Nb with 1.4 % Ni and the ADI alloyed with 0.3% Mo and 1.4 % Ni (taken as standard ADI) has been done. The results obtained for microalloyed ADIs are compared with those of base alloy having 1.4 % Ni only and the standard ADI. The following conclusions have been drawn. The tensile strength of as cast ductile iron increases remarkably with the addition of microalloying element V or Nb. The austemperability of ductile iron microalloyed with 0.1 % V or % Nb compares favorably with that of the standard ductile iron alloyed with 0.3 % Mo.

12 Paper pdf, Page 12 of 12 Microalloyed ADIs with 0.1 % V and % Nb have shown improvement in tensile strength over base ADI by 17 % and 5 %, respectively. The same over standard ADI are 13% and 2%, respectively. The improvement in mechanical properties of microalloyed ADIs are through grain refinement and precipitation hardening due to presence of the fine and uniformly distributed vanadium and niobium carbide precipitates. A reduction of 25 % and 8 % in average linear wear rates of microalloyed ADIs with 0.1 % V and % Nb, respectively, are observed over base ADI. The same over standard ADI are 16% and 5%, respectively. The increased wear resistance of microalloyed ADIs is mostly due to increase in their load bearing abilities due to the presence of hard alloy carbides precipitates in the tough ausferrite matrices. The presence of microalloying element V or Nb in ADI does not alter its basic mechanism of wear. The wear is mostly oxidative and delamination type. ACKNOWLEDGMENTS The author is extremely thankful to Prof. A Basak and Prof. A K Chakraborty of Dept. of Metallurgical and Materials Eng. IIT Kharagpur, India for their guidance in completion of this project. The author expresses his gratitude to Management of Tata Motors Jamshedpur for carrying out the experimental work. REFERENCES 1. Moore, D.J., Rouns T.N., Rundman, K.B., Structure and Mechanical Properties of Austempered Ductile Iron, AFS Transactions, vol. 93, pp (1985). 2. Rouns, T.N., Rundman, K.B., Moore, D.M., On the Structure and Properties of Austempered Ductile Cast Iron, AFS Transactions, vol. 92, pp (1984). 3. Gundlach, R.B., Janowak, J.F. Proc. 2 nd International Conference on ADI, pp.23 (1986). 4. Shah, S.M., Verhoeven, J.D., Erosion Behaviour of High Silicon Bainitic Structures I: Austempered Ductile Cast Iron, Wear, vol. 113, pp (1986). 5. Bartosiewicz, L., Krause, A.R., Alberts, F.A., Singh, I., Puttunda, S.K., Influence of Microstructure on High-Cycle Fatigue Behaviour of Austempered Ductile Cast Iron, Materials Characterization, vol. 30, pp (1993). 6. Lottridge, N.M., Grindahl, R.B., SAE, pp. 213 (1982). 7. Bahmani, M., Elliott, R., Effects of Pearlite Formation on Mechanical Properties of Austempered Ductile Iron, Materials Science and Technology, vol. 10, pp (1994). 8. Dawson, J.V., Vanadium in Cast Iron, UK International Exchange Paper, The British Foundryman, vol. 75, pp (1982). 9. Takita, M., Ueda, Y., Cast Metals, vol. 3, pp (1988). 10. Elliott, R., Cast Iron Technology, Butterworths publication (1988). 11. ASTM A897/897M-06, Standard Specification for Austempered Ductile Iron Castings. 12. Basak, A., Roy, D.K., Dutta, G.L., Wear, vol. 184, pp (1995). 13. Zhukov, A.A., Basak, A., Yanchenko, A.B., New Viewpoints and Technologies in Field of Austempering of Fe-C s, Materials Science and Technology, vol. 13, pp (1997). 14. Dubensky, W.T., Rundman, K.B., An Electron Microscope Study of Carbide Formation in Austempered Ductile Iron, AFS Transactions, vol. 93, pp (1985). 15. Dawson, J.V., Exchange Paper, Great Britain, 49 th International Foundry Congress, Chicago, Illinois (April 1982). 16. Shen, X.P., Harris, S.J., Noble, B., Influence of Small Vanadium and Cobalt Additions on Microstructure and Properties of Ductile Iron, Materials Science and Technology, vol. 11, pp (1995). 17. Lebeau, C., Production and Control of HSLA Steel Castings, AFS Transactions, vol. 92, pp (1984). 18. Skoblo, T.S., Sandler, N.I., Parfenyak, V.K., Gilman, B.S., Russian Casting Production, vol. 7, pp. 306 (1968). 19. Dorazil, E., Barta, B., Munsterova, E., Stransky, L., Huvar, A., High Strength Bainitic Ductile Cast Iron, AFS International Cast Metals Journal, pp (June 1982). 20. Ping, L., Bahadur, S., Wear, vol. 138, pp. 269 (1990). 21. Owhadi, A., Hedjazi, J. Davami, P. Wear Behavior of 1.5 Mn Austempered Ductile Iron, Materials Science and Technology, vol. 14, pp (1998). 22. Sastry, B.V.S.K., PhD Thesis on Studies on Some Characteristics of ADI, I I T Kharagpur, India, (1993). 23. Prado, J. M., Pujol, A., Cullell, J. Tarter, J., Dry Sliding Wear of Austempered Ductile Iron, Materials Science and Technology, vol. 11, pp (1995).