Primary Carbide Refinement In Hypereutectic High Chromium Cast Irons

Abstract

The invention relates to a process for manufacturing a casting of a hypereutectic high chromium cast iron (HCCI) including: forming a melt of a molten hypereutectic high chromium cast iron and a carbide modifier, and forming a solid casting of the hypereutectic high chromium cast iron. The invention further relates to a carbide-modified hypereutectic high chromium cast iron (HCCI) that includes a dispersion of a primary carbide phase in a ferrous matrix, with the primary carbide phase including a carbide modifier phase within the primary carbide phase that acted as a nucleation agent for the primary carbide phase.

Claims

1. A process for manufacturing a casting of a hypereutectic high chromium cast iron (HCCI) including: forming a melt of a molten hypereutectic high chromium cast iron and a carbide modifier, and forming a solid casting of the HCCI, wherein the carbide modifier is selected from borides of Ti, Zr, Hf, V, Nb and Ta, and wherein the melt comprises about 0.01 wt. % to 2 wt. % carbide modifier.

2. The process defined in claim 1 including forming the molten hypereutectic high chromium cast iron without the carbide modifier and adding the carbide modifier to the molten hypereutectic high chromium cast iron to form the melt.

3. The process defined in claim 2 wherein the carbide modifier is added to the molten hypereutectic high chromium cast iron in a solid form.

4. The process defined in claim 3 wherein the carbide modifier is added to the molten hypereutectic high chromium cast iron in the form of a powder.

5. The process defined in claim 3 wherein the carbide modifier is added to the molten hypereutectic high chromium cast iron in the form of a pellet.

6. The process defined in claim 2 wherein the carbide modifier is added to the molten hypereutectic high chromium cast iron in a master alloy.

7-36. (canceled)

37. The process defined in claim 1 wherein forming a solid casting of the hypereutectic high chromium cast iron includes controlling solidification of the molten hypereutectic high chromium cast iron to form equiaxed-shaped primary carbides.

38. The process defined in claim 1 wherein the carbide modifier is titanium boride and wherein the melt comprises about 0.01 wt. % to 0.8 wt. % carbide modifier.

39. A carbide-modified hypereutectic high chromium cast iron (HCCI) casting manufactured according to the process defined in claim 1, wherein the hypereutectic HCCI includes a dispersion of primary carbides in a ferrous matrix, wherein the primary carbides are at least 5% by volume smaller than primary carbides produced under the same processing conditions without the carbide modifier in the hypereutectic high chromium cast iron.

40. The carbide-modified hypereutectic high chromium cast iron defined in claim 9 wherein the hypereutectic high chromium cast iron has a lower percentage of needle-shaped primary carbides and a higher percentage of equiaxed-shaped primary carbides by volume than primary carbides produced under the same processing conditions without the carbide modifier in the hypereutectic high chromium cast iron.

41. The carbide-modified hypereutectic high chromium cast iron defined in claim 9 wherein the carbide modifier is added to comprise 0.01 wt. % to 2 wt. % of the HCCI during casting, and the carbide modifier phase is a boride of Ti, Zr, Hf, V, Nb, or Ta.

42. The carbide-modified hypereutectic high chromium cast iron defined in claim 9 wherein the carbide modifier is added to comprise 0.01 wt. % to 0.8 wt. % of the HCCI during casting, and the carbide modifier phase is titanium boride.

43. A component for use in mining and mineral processing applications manufactured from a hypereutectic high chromium cast iron defined in claim 9.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0075] The invention is described, by way of example only, with reference to the following Figures, of which:

[0076] FIG. 1 shows optical microstructures of as-cast hypereutectic HCCI samples with (a) no addition, (b) 0.03 wt. % TiB.sub.2, (c) 0.1 wt. % TiB.sub.2, (d) 0.3 wt. % TiB.sub.2, (e) 0.5 wt. % TiB.sub.2 and (f) 0.8 wt. % TiB.sub.2 at a pouring temperature of 1470? C.;

[0077] FIG. 2(a) shows a schematic illustration of microstructures of hypereutectic HCCI samples at different cross-sections;

[0078] FIG. 2(b) shows an optical microstructure of an hypereutectic HCCI sample containing 0.5 wt. % TiB.sub.2 when sectioned laterally;

[0079] FIG. 2(c) shows an optical microstructure of an hypereutectic HCCI sample containing 0.5 wt. % TiB.sub.2 when sectioned horizontally;

[0080] FIG. 3 shows optical microstructures of hypereutectic HCCI samples with (a) no addition, (b) La.sub.2O.sub.3, B containing 0.05 wt. % La.sub.2O.sub.3, (c) La.sub.2O.sub.3, B containing 0.1 wt. % La.sub.2O.sub.3, (d) La.sub.2O.sub.3, B containing 0.15 wt. % La.sub.2O.sub.3, (e) La.sub.2O.sub.3, B containing 0.2 wt. % La.sub.2O.sub.3, and (f) La.sub.2O.sub.3, B containing 0.4 wt. % La.sub.2O.sub.3 at a pouring temperature of 1470? C.;

[0081] FIG. 4 shows variations of (a) fraction of needle-shaped primary carbides; (b) the average length of needle-shaped primary carbides; and (c) the average diameter of equiaxed primary carbides, with increase in TiB.sub.2 addition in hypereutectic HCCI samples;

[0082] FIG. 5 shows optical microstructures of heat treated hypereutectic HCCI samples with (a) no addition, (b) 0.03 wt. % TiB.sub.2, (c) 0.1 wt. % TiB.sub.2, (d) 0.3 wt. % TiB.sub.2, (e) 0.5 wt. % TiB.sub.2 and (if) 0.8 wt. % TiB.sub.2;

[0083] FIG. 6(a) is a typical SEM micrograph showing the morphology of an equiaxed primary carbide containing a transformed carbide modifier phase in a hypereutectic HCCI sample with 0.5 wt. % TiB.sub.2 addition;

[0084] FIG. 6(b) is an EDS mapping showing the elementary distribution;

[0085] FIG. 7(a) is a TEM bright field image showing a primary carbide containing a transformed carbide modifier phase in a hypereutectic HCCI sample with 0.5 wt. % TiB.sub.2 addition;

[0086] FIG. 7(b) is a selected area diffraction pattern of the transformed carbide modifier phase in the hypereutectic HCCI sample;

[0087] FIG. 7(c) is a selected area diffraction pattern of the transformed carbide modifier phase in the hypereutectic HCCI sample;

[0088] FIG. 8 shows optical microstructures of as-cast hypereutectic HCCI samples with (a) 0.03 wt. % TiN, (b) 0.1 wt. % TiN, and (c) 0.2 wt. % TiN;

[0089] FIG. 9 shows variations in hardness of as cast and heat treated hypereutectic HCCI samples with increase in content of TiB.sub.2 at a pouring temperature of 1470? C.;

[0090] FIG. 10 shows variations in wear resistance of as-cast (AC) and heat treated (HT) hypereutectic HCCI samples with increase in number of cycles for no addition, 0.03 wt. % TiB.sub.2 addition and 0.1 wt. % TiB.sub.2 addition at a pouring temperature of 1470? C.;

[0091] FIG. 11 shows variations in transverse rupture strength (TRS) of as-cast and heat treated hypereutectic HCCI samples versus TiB.sub.2 concentration at a pouring temperature of 1470? C.;

[0092] FIG. 12 shows variations in TRS of as-cast heat-treated hypereutectic HCCI samples versus La.sub.2O.sub.3/B concentration at a pouring temperature of 1470? C.;

[0093] FIG. 13 shows optical microstructures of the as-cast HCCI samples with (a) no addition, (b) 0 wt. % TiN, 0.1 wt. % Ni, (c) 0.03 wt. % TiN, 0.1 wt. % Ni, and (d) 0.1 wt. % TiN, 0.1 wt. % Ni;

[0094] FIG. 14 shows optical microstructures of the as-cast HCCI samples with (a) no addition, (b) 0.03 wt. % in-situ formed TiN and (c) 0.1 wt. % in-situ formed TiN;

[0095] FIG. 15 shows variation of percent of needle-shaped primary carbides in HCCI with the addition level of TiN (a) directly added TiN and (b) in-situ formed TiN;

[0096] FIG. 16(a) is a SEM micrograph showing the directly added TiN particle in the matrix of the HCCI sample with 0.1 wt. % TiN;

[0097] FIG. 16(b) is an EDS mapping showing the elementary distribution;

[0098] FIG. 17(a) is a SEM micrograph of primary carbide containing in-situ formed TiN particle at its edge in HCCI sample with 0.1 wt. % TiN;

[0099] FIG. 17(b) is an EDS mapping showing the elemental distribution;

[0100] FIG. 18 shows variation in hardness of as-cast and heat treated HCCI sample modified with (a) directly added TiN and (b) in-situ formed TiN;

[0101] FIG. 19 shows variation in wear resistance of as-cast and heat treated HCCI samples with (a) directly added TiN and (b) in-situ formed TiN;

[0102] FIG. 20 shows variation of TRS in as-cast and heat treated HCCI samples modified with (a) directly added TiN and (b) in-situ formed TiN;

[0103] FIG. 21 shows SEM micrographs of the matrix of HCCI samples with (a) as-cast, no addition, (b) heat treated, no addition, (c) heat treated, 0.03 wt. % directly added TiN, and (d) heat treated, 0.03 wt. % in-situ formed TiN;

[0104] FIG. 22(a) is a TEM micrograph of HCCI sample with no addition showing the presence of eutectic carbides, martensite and secondary M23C6 carbides;

[0105] FIG. 22(b) shows a diffraction pattern of secondary M23C6 carbides;

[0106] FIG. 22(c) shows a diffraction pattern of martensite;

[0107] FIG. 23(a) shows an optical microstructure of the H8838A sample; and (b) shows the carbide width distribution of the H8838A sample;

[0108] FIG. 24 (a) shows an optical microstructure of the H8838B sample; and (b) shows carbide width distribution of the H8838B sample;

[0109] FIG. 25 (a) shows an optical microstructure of the H8843B sample; and (b) shows the carbide width distribution of the H8843B sample; and

[0110] FIG. 26 (a) shows an optical microstructure of the H88862A sample; and (b) shows an optical microstructure of the H88862B sample.

DETAILED DESCRIPTION OF EMBODIMENTS

Overview

[0111] Initial research and development work has established that it is possible to modify primary carbides in hypereutectic HCCIs through addition of TiB.sub.2 and La.sub.2O.sub.3/B to a hypereutectic HCCI alloy melt.

[0112] Specifically, as described further below, in the case of the addition of TiB.sub.2, the addition of 300 ppm (0.03 wt. %) TiB.sub.2 to the particular alloy melt led to a decrease in the fraction of needle-shaped primary carbides by 35%, a reduction in length of needle-shaped primary carbides by 32%, and a reduction in diameter of equiaxed carbide particles by 32%. As a result, the transverse rupture strength (TRS) of the hypereutectic HCCI alloy after heat treatment increased by 37% together with no noticeable loss in wear or hardness.

[0113] In addition, as described below, although excess TiB.sub.2 addition up to 8000 ppm (0.8 wt. %) continued reducing the amount of needle-shaped carbide, the improvement in TRS was annulled by the embrittlement of TiB.sub.2. Detailed microstructural examination showed that upon addition, TiB.sub.2 decomposed in the melt, leading to the formation of TiN particles with fresh surfaces, which acted as a M.sub.7C.sub.3 carbide refiner through promotion of the heterogenous nucleation of M.sub.7C.sub.3 carbides.

[0114] A further study investigated the effect of the form of TiN addition, particularly in the direct addition of TiN pellets and the in-situ formation of TiN. Results from this study supported the previous conclusion in that direct addition of TiN did not modify the primary carbides, while the in-situ formed TiN led to conversion of over 50% needle-shaped primary carbides into equiaxed carbides. The alloys produced with the in-situ formed TiN also exhibited similar advantages as the alloys produced by adding TiB.sub.2, in both as-cast and heat treated formed.

[0115] Further validation of results confirmed that alternative carbide modifiers could be used with similar effects, including nitrides or borides of Zr, Hf, V, Nb, and Ta. Moreover, the carbide modifiers may be added in the form of a master alloy.

Example 1Research and Development Work

Materials

[0116] The alloy melt used in the research and development work was a hypereutectic high chromium cast iron (HCCI), with chemical composition listed in Table 1.

[0117] For samples having TiB.sub.2 as a carbide modifier, 3 kg of the alloy was melted in an alumina silicate crucible in an Inductotherm induction furnace with argon as shielding gas.

[0118] TiB.sub.2 carbide modifiers were examined that were made through compressing a powder mixture of 10-30 wt. % TiB.sub.2, 1-5 wt. % Cr, 1-5 wt. % Ti and balance Fe and impurities at 150-250 MPa pressure to produce cylindrical blocks (compacts/pellets) of 15 mm in diameter and 25 mm in height. The powders used included: Fe powder with an average particle size of 75 ?m, TiB.sub.2 powder with an average particle size of less than 10 ?m, Cr powder with an average particle size of 150 ?m, and Ti powder with a particle size range of 20-63 ?m.

[0119] Further La.sub.2O.sub.3/B carbide modifiers were also examined, comprising 8-32% La.sub.2O.sub.3, 3-12 wt. % B, 1-5 wt. % Cr, 7-13 wt. % Ni and balance Fe and impurities, which were produced by a similar method as the TiB.sub.2 carbide modifiers.

[0120] Modification of a hypereutectic HCCI alloy was achieved by dropping 10 g compacts/pellets containing TiB.sub.2 or La.sub.2O.sub.3/B as carbide modifiers into a melt of a HCCI alloy. Alternatively, the powder mix or compacts/pellets could be melted to form a master alloy, for example, in a Plasma Arc Furnace (PAM) under Argon atmosphere.

[0121] The compacts/pellets were pre-heated to 600? C. before introducing the compacts/pellets into the melt followed by electromagnetic stirring for 10 minutes.

[0122] This procedure enabled different TiB.sub.2 and La.sub.2O.sub.3, B additions ranging from 0.03 wt. % to 0.8 wt. %.

[0123] In order to understand the effect of carbide modifiers at different pouring temperatures, the temperature of the melt was measured before pouring using thermometer fitted with R-type contact block of lance pipe. The other end of the lance pipe was fitted with multi-dip thermocouples whose tip is dipped in the melt to measure the temperature before pouring. After the melt reached the desired range of pouring temperature, the melt was poured in CO.sub.2-silicate moulds, thereby forming cast samples.

[0124] The dimensions of the samples obtained from sand casting were 55 mm?20 mm?10 mm.

TABLE-US-00001 TABLE 1 Chemical composition of the base hypereutectic high chromium cast iron Cr C Mn Si Mo Ni Cu S P Fe 25-35 2-8 1-5 ?2 ?2 ?2 ?0.5 ?0.5 ?0.5 Bal.

[0125] After casting, half of the samples for each TiB.sub.2 addition were heat treated at 1000? C. for 2 hours followed by air cooling. This heat treatment is a standard destabilization process for this type of HCCIs.

Microstructural Characterization

[0126] A section of each of the as-cast samples obtained was cut from the centre.

[0127] The cutting was done using Struers 66A25 cutting blades. The section was then hot mounted in polyfast resin and mechanically grounded, polished and then etched with ferric chloride solution.

[0128] For preparing samples for metallographic observation, the mounted samples were first ground on silicon carbide sand paper up to 120 grit size. Then the samples were polished on Largo, Mol, and Nap respectively. Samples were then etched with ferric chloride solution, washed with ethanol and dried with compressed air before examining it under the Reichert-Jung Polyvar Met optical microscope and JEOL 6610 Scanning Electron Microscope.

[0129] X-Ray Diffraction (XRD) analysis of the samples was done using Bruker D8 Advanced Powder X-Ray diffractometer. The samples were scanned at 2? angle between 35-90 with a step size and scan rate were 0.02? and 0.5?/min respectively.

[0130] FEI Scios Focused Ion Beam (FIB)dual beam SEM system was used to prepare TEM samples by slicing the carbide containing the modifier and mounting it on a copper grid.

[0131] A Hitachi HF5000 transmission electron microscope operated at an accelerating voltage of 200 kV was used to obtain the selected area diffraction pattern of the primary carbides and carbide modifiers.

Mechanical Property Tests

[0132] The hardness of all hypereutectic HCCI samples was measured with HRC (Hardness Rockwell C).

[0133] Wear tests were conducted on sand paper using a Taber Rotary Abraser 5135. Samples with cross-section 10 mm?10 mm were cut from the as-cast and heat-treated samples and ground to obtain a flat surface. The flat surfaces of samples were placed on the sandpaper disc of the abraser with a load of 1.5 kg. 120 grit SiC sand paper was used as abrasive material. The rotation speed of the disc was set at 72 rpm. After every 200 rotations, the sandpaper was changed.

[0134] The weight of each sample was recorded before and after every 1000 rotations. The tests were up to 5000 rotations.

[0135] The variation of weight loss of each sample with the number of rotations was used to measure the wear resistance of the hypereutectic HCCI alloy.

[0136] To evaluate the influence of the carbide modification on the fracture toughness of the HCCIs, transverse rupture strength (TRS) or bending strength was determined using samples of 10 mm?10 mm?50 mm that were also cut from the as-cast and heat-treated samples.

[0137] The three-point bend test was done in an Instron 4505 machine along with Bluehill 2 software, according to the ASTM standard B528-16 [25]. The load was applied at a constant rate of 5 mm/min until fracture. The TRS value was calculated using the equation below.

TRS=(3?P?L)/(2?T{circumflex over ()}2?W)

where P is the force measured to fracture the specimen, L is the distance between the centers of the supporting rods, T and W are the thickness and the width of the specimen respectively.

[0138] The value of L as per standard was fixed at 25.4 mm. The thickness and width of the specimen were 10 mm?10 mm.

[0139] For each sample type, four specimens were tested and the average of the four TRS values was used for the plotting.

Results and Discussion

Microstructural Characterization

[0140] FIG. 1 shows microstructural evolution of the primary Cr.sub.7C.sub.3 carbide with increase in TiB.sub.2 addition in as-cast HCCI samples, from (a) no addition, (b) 0.03 wt. % TiB.sub.2, (c) 0.1 wt. % TiB.sub.2, (d) 0.3 wt. % TiB.sub.2, (e) 0.5 wt. % TiB.sub.2 and (f) 0.8 wt. % TiB.sub.2 at the pouring temperature of 1470? C.

[0141] In the sample without TiB.sub.2 as shown in FIG. 1(a), the microstructure consists of needle-shaped and equiaxed primary carbide particles and eutectic structure composed of the same M.sub.7C.sub.3 carbide and ferrous matrix. Part of the matrix was transformed to pearlite/martensite/bainite depending on cooling rate together with large amount of retained austenite. As the research and development focused on modification of the primary carbides, details of the eutectic structure are not described.

[0142] In order to clarify the influence of sample orientation on the observed morphology of the primary carbide, metallographic examination was done on three different cross sections of the sample.

[0143] FIG. 2(a) is a schematic illustration of primary carbides on three different cross-sections, while FIGS. 2(b) and (c) are optical microstructures of HCCI containing 0.5 wt. % TiB.sub.2 when sectioned (b) laterally and (c) horizontally.

[0144] As observed in FIGS. 2(a), 2(b) and (c), the microstructural morphology of the as-cast sample containing 0.5 wt. % TiB.sub.2 is almost identical on different cross sections. This implies that the small carbide particles are equiaxed. This also demonstrates that sampling from different orientations of the sample had no effect on the observed microstructure.

[0145] In addition, at higher magnification as shown in the insert of FIG. 1(a), cracks and pores are visible in both the needle and equiaxed carbide particles, indicating the high brittleness of the carbides.

[0146] Addition of 0.03 wt. % TiB.sub.2 dramatically modified the morphology of the primary carbides as shown in FIG. 1(b). Not only did the volume fraction of needle carbide significantly decrease, but the average size of both needle-shaped and equiaxed carbide particles was also reduced.

[0147] However, further addition of TiB.sub.2, i.e., 0.5 wt. % or higher TiB.sub.2, led to very marginal change in both the fraction of the needle carbide and the size of the equiaxed carbide particles as shown in FIG. 1(c)-(e).

[0148] Increasing the addition of TiB.sub.2 to 0.8 wt. % resulted in slightly coarsening of the equiaxed carbide particles as shown in FIG. 1(f).

[0149] To quantitatively investigate the influence of TiB.sub.2 addition on the morphology of the primary carbide, variations of volume fraction of the needle-shaped carbides (relative to the total microstructure), the average length of the needle-shaped carbide particles and the average diameter of the equiaxed carbide particles with the TiB.sub.2 addition are plotted in FIG. 4.

[0150] Although the average width of the needle primary carbides did not vary much with the addition, the volume fraction and the average length of the needle carbides significantly reduced from 10.8% and 686.37?55.51 ?m, respectively, in the sample without TiB.sub.2 down to 6.98% and 462.71?23.69 ?m, respectively, in the sample with 0.03 wt % TiB.sub.2. Furthermore, the average size of the equiaxed primary carbides also reduced from 87.15?2.09 ?m to 59.27?1.21 ?m.

[0151] The increase in volume fraction and reduction in size of the equiaxed primary carbides were attributed to the TiB.sub.2 particles that promoted the heterogenous nucleation of the primary carbides.

[0152] The high nucleation rate also constrained the growth of the carbide particles along one particular direction to form needle-shaped carbides, reducing the volume fraction of needle carbides. The mechanism of carbide modification by TiB.sub.2 is discussed further below.

[0153] The slight coarsening of the equiaxed carbide particles at high content of TiB.sub.2 addition as shown in FIG. 1(f) is related to the so called poisoning effect. The actual cause of poisoning effect is unclear. It may be associated with the formation of constitutional supercooling (CS) zone at front of the growing solid/liquid interface as a result of rejection of solutes from the solid (in the case of M.sub.7C.sub.3 primary carbide, the solute could be boron resulted from the decomposition of TiB.sub.2). This CS zone may promote the growth of the solid if the CS zone was not big enough to activate additional heterogenous nucleation within the CS zone or there was lack of proper nucleant within the zone.

[0154] FIG. 3 shows the effect of variation of La.sub.2O.sub.3, B on the primary carbides in hypereutectic HCCI samples at a pouring temperature of 1470? C., specifically at (a) no addition, (b) 0.05 wt. % (La.sub.2O.sub.3, B), (c) 0.1 wt. % (La.sub.2O.sub.3, B), (d) 0.15 wt. % (La.sub.2O.sub.3, B), (e) 0.2 wt. % (La.sub.2O.sub.3, B), and (f) 0.4 wt. % (La.sub.2O.sub.3, B). As the content of La.sub.2O.sub.3/B increased, there was a decrease in size of primary carbides up to 0.2 wt. %. Any further increase in of La.sub.2O.sub.3/B content not only increased the size of the primary carbides but also the numbers of internal cracks in the samples, as shown in FIG. 3(f). Due to chemical similarities, it is believed that similar advantages may be achieved with a Ce.sub.2O.sub.3/B carbide modifier.

[0155] Generally, as-cast hypereutectic HCCIs are very rarely used because the un-tempered martensite/bainite and large amount of stable retained austenite. Destabilization treatment is commonly performed on HCCI components.

[0156] FIG. 5 shows the microstructure of HCCI samples after heat treatment at 1000? C. for 2 hours for (a) no addition, (b) 0.03 wt. % TiB.sub.2, (c) 0.1 wt. % TiB.sub.2, (d) 0.3 wt. % TiB.sub.2, (e) 0.5 wt. % TiB.sub.2 and (f) 0.8 wt. % TiB.sub.2. The Figure shows that the heat treatment did not change the morphology of the primary carbides, while the eutectic matrix changed.

[0157] Generally, the heat treatment resulted in transformation of the matrix from austenite to martensite and retained austenite. While the eutectic M.sub.7C.sub.3 carbides remain unaffected by the destabilization treatment, there is precipitation of M.sub.23C.sub.6 another secondary carbide in the matrix.

Modification Mechanism of the Primary Carbide by TiB.SUB.2

[0158] In order to understand how the TiB.sub.2 addition modified the primary Cr.sub.7C.sub.3 carbide, microstructures of the as-cast hypereutectic HCCI samples were further examined in electron microscopes.

[0159] FIG. 6(a) shows a typical SEM micrograph of a primary carbide containing what is described herein as an inoculant particle in the HCCI sample with 0.5 wt. % TiB.sub.2 addition (a higher TiB.sub.2 addition level was used in order to easily identify the inoculant within the carbide).

[0160] In FIG. 6(b), EDS mapping analysis indicated that the carbides is rich in Cr and Mn, but the inoculant particle is rich in titanium and nitrogen. Hence, although the primary carbide was confirmed to be a Cr carbide, EDS mapping analysis indicated that the inoculant particle (phase) was TiN.

[0161] This type of inoculant particles was observed in most of the equiaxed primary carbide particles in all the hypereutectic HCCI samples with TiB.sub.2 addition.

[0162] TEM examination was conducted to further characterize the inoculant particle.

[0163] FIG. 7(a) shows a typical TEM micrograph of a primary carbide containing an inoculant in the hypereutectic HCCI sample with 0.5 wt. % TiB.sub.2 addition. The selected area diffraction patterns in FIG. 7(b) for diffraction from the inoculant particle and FIG. 7(c) for diffraction from the primary carbide verified that the primary carbide is hexagonal Cr.sub.7C.sub.3 and the crystal structure of the inoculant particle is the same as TiN.

[0164] From above results, it follows that the inoculant that promoted the heterogenous nucleation of primary M.sub.7C.sub.3 carbide in the hypereutectic HCCI samples was TiN.

[0165] Over 120 pairs of TiN nuclei and M.sub.7C.sub.3 primary carbides were examined with EBSD analysis using the method proposed by Qiu and co-workers (Qiu, D., M.-X. Zhang, and P. M. Kelly, Crystallography of heterogeneous nucleation of Mg grains on Al2Y nucleation particles in an Mg-10 wt. % Y alloy. Scripta Materialia, 2009. 61(3): p. 312-315 DOI: 10.1016/j.scriptamat.2009.04.011.).

[0166] However, no reproducible orientation relationships (ORs) between TiN and M.sub.7C.sub.3 carbide in the HCCI samples were determined.

[0167] In order to further understand the effect of TiN, commercial 0.03-0.2 wt. % TiN powder with average particles size of less than 10 ?m was added into the hypereutectic HCCI in the same way as adding TiB.sub.2.

[0168] As shown in FIG. 8, the efficiency of the externally added TiN particles on modification of the primary carbide of the hypereutectic HCCI is very marginal for (a) 0.03 wt. % TiN, (b) 0.1 wt. % TiN, and (c) 0.2 wt. % TiN.

[0169] This implies that externally added TiN particles are not good modifier for M.sub.7C.sub.3 primary carbide. The oxide thin film on the particles may have been responsible for such low efficiency as the oxide film could influence the wetting of the melt on the particles.

[0170] Hence, in situ formed fresh TiN particles with clean surfaces are necessary to promote the heterogenous nucleation of Cr.sub.7C.sub.3 in the melt. Currently, research work is being conducted to verify this assumption.

Mechanical Properties

Hardness and Wear Resistance

[0171] FIG. 9 shows the variation of hardness of hypereutectic HCCI with addition of TiB.sub.2 before and after heat treatment.

[0172] Although the heat treatment resulted in a significant increase in hardness due to more martensite formed in the matrix as a result of destabilization treatment, the effect of TiB.sub.2 addition on hardness is very marginal even though slightly increase in hardness can be observed at 0.03 wt. % addition.

[0173] Properties of hypereutectic HCCI are essentially governed by the primary carbide and microstructure of the eutectic matrix.

[0174] Although TiB.sub.2 addition refined the primary carbide, there was no influence on the eutectic matrix. In addition, additions over 0.1 wt. % also increased the cracks and pores in the carbide and matrix. Carbide refinement enhanced the uniformity and dispersion of the carbide particles, which increased hardness. Thus, at 0.03 wt. % TiB.sub.2 addition, the hardness was slightly increased. But, cracks and pores at higher addition levels could also negatively affect the properties, which sets off the effect of the refinement.

[0175] Hence, the overall influence of TiB.sub.2 addition on hardness of the hypereutectic HCCI samples is marginal. The reduction in hardness of the as-cast sample with 0.1 wt. % TiB.sub.2 addition could be attributed to the high cast defect as the data was more scatter.

[0176] The influence of TiB.sub.2 addition on wear resistance of hypereutectic HCCI samples is consistent with the hardness.

[0177] FIG. 10 shows the wear loss of hypereutectic HCCI samples with 0 wt. %, 0.03 wt. % and 0.1 wt. % TiB.sub.2 additions before and after heat treatment at different wear cycles on sand papers.

[0178] The destabilization treatment effectively improved the wear resistance of samples at all TiB.sub.2 addition levels as a result of formation of more martensite in the matrix. Although the carbide modification has almost no effect on wear performance of as-cast alloys, addition of 0.03 wt. % TiB.sub.2 slightly increased the wear resistance after heat treatment at high wear cycles. This is related to the more dispersal distribution of primary carbide particles as a result of carbide refinement.

[0179] At high TiB.sub.2 addition level, the higher density of cracks and pores setoff the effect of the carbide refinement, leading to similar wear behavior of the modified alloy as that of the un-modified alloy.

Transverse Rupture Strength (TRS)

[0180] Fracture toughness of the HCCI with and without addition of TiB.sub.2 was measured using 3-point bending tests.

[0181] FIG. 11 shows the variation of TRS with TiB.sub.2 addition at as-cast and heat-treated conditions.

[0182] At as-cast condition, the TRS values remains unchanged with minor addition of 300 ppm TiB.sub.2.

[0183] However, after destabilisation treatment, the 300 ppm TiB.sub.2 addition led to a significant increase in TRS value by more than 38% compared the alloy without TiB.sub.2 addition. The destabilisation treatment led to the transformation of austenite in the eutectic matrix to martensite. In this case, the modification of primary carbide is more effective in preventing crack propagation, increasing the fracture toughness.

[0184] Further addition of TiB.sub.2 decreased the TRS value at both the as-cast and heat-treated conditions due to the formation of too much TiN particles as a result of decomposition of TiB.sub.2 in the melt.

[0185] The variation in TRS value with increase in La.sub.2C.sub.3, B in the alloy is shown in FIG. 12. The Figure shows that there was a negative effect of La.sub.2C.sub.3, B content on the TRS value of as-cast HCCI. Upon heat treatment at 1000? C. for 2 hours, there was an increase in TRS with a maximum at 0.15 wt. %. Any more additions of La.sub.2C.sub.3/B in the alloy resulted in a decrease in TRS value.

CONCLUSION

[0186] Addition of TiB.sub.2 particles with average size of 10 ?m is an effective method to modify and refine primary carbides in hypereutectic HCCIs.

[0187] With an optimal addition of 300 ppm (0.03 wt. %) in the hypereutectic HCCI evaluated in the research and development work, the fraction of needle-shaped primary carbides was reduced by 35%, a reduction in length of the needle-shaped by 32% and a reduction in diameter of the equiaxed carbide particles by 32%.

[0188] The work found that TiB.sub.2, upon addition into the hypereutectic HCCI melt, decomposed into Ti and B, leading to formation of TiN particles, which further acted as heterogenous nucleation sites for the solidification of the primary carbide. This promoted primary carbide nucleation, decreasing the fraction and length of needle-shaped carbide and refining both the equiaxed and needle-shaped carbide particles in the alloy. This was apparent for each as-cast and heat treated sample with TiB.sub.2 addition compared to the, compared to the samples with no TiB.sub.2 addition.

[0189] Furthermore, addition of 300 ppm (0.03 wt. %) TiB.sub.2 increased the TRS value of the destabilized hypereutectic HCCI by more than 37% without reduction in both hardness and wear resistance compared with the unmodified alloy as the result of primary carbide modification.

[0190] But, further additions led to decrease in TRS because of the embrittlement caused by the excess amount of TiN formed.

[0191] When TiB.sub.2 carbide modifier was introduced into HCCI melt, it reacted with contents of the melt to form TiN, which can act as an effective modifier for the primary carbides

[0192] Due to the low additions of TiB.sub.2, the volume fraction of primary carbides remains almost unchanged. Therefore, there was slight or no variation in properties such as hardness or wear resistance.

[0193] In the as-cast condition, there was little or no change in TRS value with increase in TiB.sub.2 content at both high and low pouring temperatures.

[0194] La.sub.2O.sub.3/B also acted as an effective carbide modifier.

Example 2Effect of TiN Addition Form

[0195] In Example 1, TiB.sub.2 was identified as an effective modifier for modification of primary carbides. It was found that, upon inoculation, the TiB.sub.2 decomposes and forms TiN, which acts as nucleation site for primary carbides and refines the primary carbides in the final casting.

[0196] As TiB.sub.2 addition also introduces boron into the alloy (which may unwanted for certain alloys), alternative forms of carbide modifier addition were investigated.

[0197] Where possible, tests were conducted in a similar manner as Example 1 for ease of comparison.

Materials

[0198] The alloy melt used in this Example was a commercial hypereutectic high chromium cast iron (HCCI) with a composition in accordance with Table 1 and prepared in accordance with Example 1.

[0199] Analysis of certain alloy melts comprised about Fe-31.1Cr-4.31C, and were in accordance with ISO 21988 Grade Cr35.

[0200] For each sample, 3 kg of the alloy was melted in an alumina silicate crucible in an Inductotherm induction furnace with argon as shielding gas. A TiN carbide modifier was added to this melt in various forms.

[0201] In particular, a TiN carbide modifier compact/pellet was produced in a similar manner as in Example 1, produced by compressing a powder mixture of about 10 wt. % TiN, 2 wt. % Ni and balance Fe and impurities to form a pellet of 14 mm in diameter and 10 mm in height. The powders used included: Fe powder with an average particle size of about 100 ?m, TiN powder with an average particle size of about 50 ?m, and Ni powder with an average particle size of about 50 ?m.

[0202] The carbide modifier pellets were preheated at 600? C. for 1 hour before directly introducing it into the melt with TiN addition levels of 0 wt. %, 0.03 wt. % and 0.1 wt. %.

[0203] The direct addition of the TiN pellet carbide modifier was compared to an in-situ formed TiN produced by melting a commercially available pure Ti block and a nitrided low carbon steel with N content of about 11%, which were added into the HCCI melt in a stoichiometric ratio of Ti to N. In the samples, a full reaction was assumed for the added Ti and Ni in the melt, and additions of 0.03 wt. % and 0.1 wt. % TiN were targeted.

[0204] After addition of either the pellet or the Ti-steel block, the melt was electromagnetically stirred for 10 minutes. The melt temperature was checked using a thermometer fitted with R-type contact block of lance pipe. Once the melt reached the desired pouring temperature of 1470? C., it was poured into CO.sub.2-silicate sand moulds of dimensions 55 mm?20 mm?10 mm.

[0205] After casting, half of the ingots for each addition was heat treated at 1000? C. for 2 hours followed by air cooling. This heat treatment is a standard destabilization process for this type of HCCIs.

Microstructural Characterization

[0206] Metallographic samples were cut from the centre of both the as-cast and heat treated ingots for each addition form using Struers 66A25 cutting blades. After hot mounting in polyfast resin, samples were mechanically grounded and polished, followed by etching with ferric chloride solution. Microstructural examination was done under a Reichert-Jung Polyvar Met optical microscope and in a Hitachi SU 3500A Scanning Electron Microscope.

[0207] FEI Scios focussed ion beam (FIB) was used to prepare transmission electron microscope (TEM) thin film and was examined in a Hitachi HF 5000 TEM operating at an accelerating voltage of 200 kV.

Mechanical Property Tests

[0208] Data on hardness, wear, weight variations, and TRS were obtained on as-cast and heat treated samples in a manner similar to that described in Example 1.

Results and Discussion

Microstructural Characterization

[0209] FIG. 13 shows the microstructure of the HCCI with directly added TiN for (a) no addition, (b) 0 wt. % TiN, 0.1 wt. % Ni, (c) 0.03 wt. % TiN, 0.1 wt. % Ni, and (d) 0.1 wt. % TiN, 0.1 wt. % Ni.

[0210] As in FIG. 13(a), the sample with no addition consists of needle-shaped primary M.sub.7C.sub.3 carbides and eutectic structure. The eutectic structure is comprised of carbide and austenite/martensite.

[0211] As shown in FIG. 13(b), addition of 0.1 wt. % Ni without TiN does not affect the primary carbides as it could dissolve in austenite as a ?-stabilizer. Increasing the content of TiN to 0.03 wt. % and 0.1 wt. % did not produce any significant effect on primary carbides as in shown in FIGS. 13(c) and 13(d).

[0212] FIG. 14 shows microstructures of the HCCI ingots with in-situ formed TiN for (a) no addition, (b) 0.03 wt. % in-situ formed TiN and (c) 0.1 wt. % in-situ formed TiN. The in-situ formation of 0.03 wt. % TiN led to a significant decrease in the size of the primary carbides as shown in FIG. 14(b). However, further increase the in-situ formed TiN to 0.1 wt. % coarsened primary carbides as shown in FIG. 14(c).

[0213] FIG. 15 quantitively demonstrates the effect of direct TiN addition compared to in-situ formed TiN on the variation of volume fraction of needle-shaped primary carbides in the resultant alloy ingots. As shown, direct addition of TiN has no significant impact on the needle-shaped primary carbides as in FIG. 15(a). As discussed in Example 1, this is likely due to surface contamination/oxidization of the TiN particles inhibiting nucleation of the primary M.sub.7C.sub.3 carbides.

[0214] In contrast, as shown in FIG. 15(b), with 0.03 wt. % of in-situ formed TiN, the percent of needle-shaped primary carbides reduced significantly from 15.94% to 6.35%. The decrease in fraction of needle-shaped primary carbides can be attributed to the in-situ formed TiN particles which act as the nucleation site for the heterogenous nucleation of primary carbides.

[0215] The mechanism of modification of primary carbides by in-situ formed TiN but not by directly added TiN will be discussed in detail in the next section.

[0216] The increase in needle shaped carbides at 0.1 wt. % of in-situ formed TiN as shown in FIG. 14(c) and FIG. 15(b) may be related to a so called poisoning effect on grain refinement observed in casting of light alloys (e.g. aluminium-based alloys) and ferritic steels; however, the mechanism of this effect is not known to the inventors. The cracks observed in FIG. 14(c) may also be attributed to this effect, combined with increase in brittleness of modified HCCI.

[0217] Example 1 previously identified similar effects upon excess addition of TiB.sub.2 (which transforms to TiN in the melt) and led to increase in needle-shaped carbides and coarsening of equiaxed carbides.

[0218] Accordingly, as it appears that adding too much in-situ formed TiN might lead to segregation of such particles along boundaries of eutectic structures as they dissolved in neither the eutectic carbide nor the eutectic austenite in the subsequent eutectic reaction after the formation of the primary carbide, which eventually led to increase in brittleness of the alloy. Therefore, the addition level of Ti and N should be optimized when the in-situ formed TiN is used to modify the primary carbide in HCCIs. In the present work, the critical value to effectively modify primary carbides is 0.03 wt. %; however it will be appreciated that this critical value may differ for different HCCI compositions or manufacturing conditions. In particular, for certain HCCI compositions or manufacturing conditions, it is believed that additions of up to 2 wt. % may be beneficial.

[0219] Similar to Example 1, after casting half of the samples were heat treated at 1000? C. for 2 hours. This heat treatment has no effect on the morphologies and fraction of the primary and eutectic M.sub.7C.sub.3 carbides in the castings, but have effects on the microstructure of the metallic matrix.

Modification Mechanism from Direct Addition TiN and In-Situ Formed TiN

[0220] To further understand the effect of TiN addition form of only the in-situ formed TiN could modifying the primary carbides (but not the directly added TiN), further examination of both the samples were done under the scanning electron microscope (SEM).

[0221] FIG. 16 shows a SEM micrograph of the HCCI with directly added 0.1 wt. % TiN particles, showing a TiN particle in the matrix. A higher TiN addition was chosen to conduct SEM examination for easier identification of the particles. FIG. 16(b) shows the EDS mapping which clearly indicates that the inoculant particle is enriched with Ti and N (contrast of N mapping is low due to the low atomic number), confirming the addition of TiN.

[0222] As TiN is a thermodynamically stable phase and has the melting point of 2930? C., it is highly unlikely to decompose or react to form other compounds. The inefficiency of externally added TiN particles in modification of the primary carbides may be related to its morphology. From FIG. 16(a), it can be seen that the TiN is an irregular shape and consists of a few smaller particles. This implies that the agglomeration occurred in the directly added TiN particles, which led to no modification of primary carbides even though TiN was added into the HCCI melt.

[0223] In comparison, FIG. 17(a) shows a typical SEM micrograph of the HCCI with 0.1 wt. % in-situ formed TiN particles. A rectangle shaped particle can be observed at the edge of a primary carbide in FIG. 17(a). The EDS elementary mapping in FIG. 17 (b) confirms enrichment in both Ti and N. As Ti has more affinity to N than to C, TiN is thermodynamically stable and is the preferred phase, confirming the in-situ formation of TiN particles in the HCCI melt.

[0224] From the above results, it can be concluded that directly added TiN did not modify primary carbides while in-situ formed TiN led to heterogenous nucleation of primary carbides causing its modification.

[0225] In the case of directly added TiN, the pellet containing TiN powder was preheated at 600? C. for 1 hour to remove moisture and air trapped in the pellet; however, preheating might cause TiN oxidation, leading to formation of oxide layer on TiN particles. Due to the absence of reactive elements such as Ti, it is believed that this oxide layer could not be broken, resulting in adverse impact on the wetting of TiN particles by HCCI melt. Furthermore, low wettability of the oxide layer in the HCCI melt, combined with difference in density between HCCI melt (7.3) and TiN (5.4), may have led to the rejection of TiN particles in to the HCCI melt. In addition, the agglomeration of TiN particles also reduced the efficiency the TiN particles to serve as inoculant for the primary carbides. Therefore, there was no effect of directly added TiN on the modification of primary carbides.

[0226] In contrast, when Ti and N are added simultaneously to the HCCI melt, there occurs in-situ formation of TiN as it is thermodynamically more stable than other phases, such as TiC. The HCCI melt contains more than 30% chromium which is much higher than the optimum content of chromium required to fully wet the freshly formed surfaces of TiN particles. This wetting helps TiN particles to promote heterogenous nucleation of primary carbides.

Mechanical Properties

Hardness and Wear Resistance

[0227] FIG. 18 shows the variation in hardness of HCCI samples modified with directly added and in-situ formed TiN in both as-cast and heat treated alloys.

[0228] As discussed previously, properties such as hardness and wear resistance in hypereutectic high chromium cast iron are governed by primary and eutectic carbides and microstructure of the matrix. Upon destabilization heat treatment, the primary and eutectic carbides remain unaffected while the matrix transforms from austenite to martensite. This change in the matrix leads to the increase in overall hardness of the HCCI samples with or without modification. However, as directly added TiN has no effect the microstructure, while in-situ formed TiN only modifies the size of primary carbides but not the volume fraction, the hardness in both cases remains relatively unchanged with increase in TiN content.

[0229] FIG. 19 shows the variation in wear resistance of HCCI samples with directly added and in-situ formed TiN particles, before and after heat treatment at different cycles on sand paper. As shown, the destabilization heat treatment did not have a significant impact on the wear performance of HCCI samples with directly added TiN and in-situ formed TiN.

Transverse Rupture Strength (TRS)

[0230] The main purpose of carbide modification and performing destabilization heat treatment is to reduce the brittleness of the HCCI. As HCCI is a typical brittle alloy, its fracture toughness is evaluated using three-point bending tests in the present work.

[0231] FIG. 20 shows the variation in TRS of the as-cast and heat treated HCCI samples with directly added and in-situ formed TiN. For the as-cast HCCI, direct addition of TiN did not significantly impact the TRS at 0.03 wt. % addition, but reduced the TRS at 0.1 wt. % addition. Destabilization heat treatment did not change the TRS of the HCCI samples without TiN addition, but the TRS value was increased by ?20% for the HCCI samples with directly added TiN after the destabilization as shown in FIG. 20(a).

[0232] With the as-cast HCCI samples with in-situ formed TiN, the variation of TRS with increasing TiN content is very marginal and is not significant. However, the destabilization heat treatment led to a ?29% increase in the TRS of the HCCI sample with 0.03 wt. % in-situ formed TiN.

[0233] Fracture toughness of the hypereutectic HCCI relies both the morphology of primary carbide and microstructure of the matrix. The latter is generally includes eutectic ledeburite. However, modification of primary carbide alone cannot improve the toughness of the as-cast HCCI because the matrix predominates the crack propagation. The eutectic matrix consisting of carbide and austenite/martensite is the preferred path of the crack due to its relative low strength and high brittleness. Therefore, heat treatment aims to transform the majority eutectic austenite to martensite to increase the strength of the matrix. In this case, morphology of primary carbide plays roles in crack propagation. Hence, modifying the primary carbide can lead to improvement of the fracture toughness. Therefore, it is preferred that, in order to obtain higher fracture toughness of HCCIs, both primary carbide modification and heat treatment are conducted based on the HCCI composition.

[0234] The increase in TRS after heat treatment was investigated in inspection of the matrix, as shown in FIG. 21, showing SEM micrographs of the matrix of HCCI samples with and without TiN additions. FIG. 21(a) shows the eutectic matrix of the as-cast HCCI sample. Destabilization treatment led to precipitation of secondary carbides and transformation of the eutectic austenite to martensite each in the HCCI samples, as shown in FIG. 21(b) for no TiN addition, FIG. 21(c) for 0.03 wt. % direct addition TiN, and FIG. 21(d) for 0.03 wt. % in-situ formed TiN.

[0235] The form of secondary carbides and matrix in the heat treated HCCI sample with no TiN addition was examined visually through use of a transmission electron microscope (TEM). FIG. 22(a) shows the typical TEM micrograph consisting of eutectic carbides along with secondary M.sub.23C.sub.6 carbides in martensitic matrix. The selected area diffraction patterns shown in FIG. 22(b), for the secondary M.sub.23C.sub.6 carbides, and (c), for the martensite, verify that the secondary carbide is FCC M.sub.23C.sub.6 carbide and the matrix is martensite. Hence, the highest TRS value obtained in the HCCI sample modified with 0.03 wt. % in-situ formed TiN was a result of combination of carbide modification and precipitation of secondary carbides and transformation of austenite to martensite during heat treatment, which increased the strength of the eutectic matrix and refined the primary carbides.

CONCLUSION

[0236] The work found that directly added TiN did not modify the primary carbides in the HCCI samples, while in-situ formation of 0.03 wt. % TiN led to conversion of over 50% needle-shaped primary carbides into equiaxed carbides.

[0237] This is likely due to the in-situ formed TiN particles having a fresh surface acting as nucleation sites promoting the heterogenous nucleation of primary carbides in HCCI, leading to modification and refinement of the carbides.

[0238] Results found that addition of 0.03 wt. % in-situ formed TiN in HCCI increased the fracture toughness (TRS) of the HCCI by about 29% after heat treatment, while direct addition of 0.03 wt. % TiN in HCCI increases TRS by 23% after heat treatment.

[0239] Both direct addition and in-situ formation of TiN did not change the hardness and wear resistance of the as-cast and as-destabilization treated HCCI. Destabilization treatment of the as-cast HCCI enabled precipitation of secondary M.sub.23C.sub.6 carbides together with the transformation of eutectic austenite to martensite increasing the strength, which contributes to the increase in TRS.

Example 3Further Carbide Modifiers

[0240] This example sought to validate the results of Examples 1 and 2 and confirm if alternative proposed carbide modifiers would offer similar advantages in reducing the fraction of needle-shaped primary carbides.

Materials

[0241] The base alloy melt used in this Example was a hypereutectic high chromium cast iron (HCCI) with a composition in accordance with Table 1. Specifically, the alloy composition was targeted to include 30 wt. % C, 4-4.5 wt. % C, variable carbide modifiers, and balance Fe and impurities. The tested carbide modifiers included ZrB.sub.2 and TiB.sub.2 at 0-0.5 wt. % addition.

[0242] The carbide modifiers were prepared as master alloys for addition into the base alloy melt to promote dispersion of the carbide modifier, while also reducing any carbide modifier losses during addition, controlling compositions ranges, and minimizing temperature differentials at addition when compared to direct addition of carbide modifier pellets or solid carbide modifier components for in-situ formation of the carbide modifier. The master alloys for ZrB.sub.2 and TiB.sub.2 were iron-based alloys.

[0243] The resultant melt was then poured into moulds, with a targeted pouring temperature of 1422? C., and air cooled.

Microstructural Characterization

[0244] Metallographic samples were cut from each of the as-cast samples and examined with an optical microscope for the characterising the formed primary carbides. In particular, the distribution of primary carbide sizes was noted; however in some samples the number of primary carbides, % area of primary carbides, width of the largest primary carbide, average width of the 20 largest primary carbides, and average length/width (L/W) ratio of the 20 largest primary carbides were noted. The width of the primary carbides were measured as the longest distance on the minor axis of the carbide.

Results and Discussion

Microstructural Characterization

[0245] FIG. 23(a) shows the microstructure of sample H8838A with a composition of Fe-30C-4.4C and no carbide modifier addition. The sample was poured at 1422? C. As shown in FIG. 23(b), the resultant alloy had a high variance of carbide width with a median width of about 50-60 ?m.

[0246] Further subsamples were cut from the H8838A alloy and the microstructure of subsample was characterized, as summarized in Table 2 below.

TABLE-US-00002 TABLE 2 Samples without carbide modifier addition AVG AVG Width of width of L/W largest top 20 ratio of No of % area of primary primary top 20 primary primary carbide carbides primary Sample carbides carbides (?m) (?m) carbides H8838A 162 18.31 164 113 1.65 H8838A 0004 173 18.09 138 107 1.52 H8838A 0008 150 20.65 209 106 2.17 H8838A 0013 116 21.03 152 102 2.99 H8838A 0014 161 18.92 130 102 1.89 AVERAGE 152 19.40 158 106 2.05

[0247] Further samples with zinc-based and titanium-based carbide modifiers were prepared and compared to the H8838A control.

Zinc-Based Carbide Modifiers

[0248] In view of the advantages provided by TiB.sub.2 addition converting in-situ to a TiN carbide modifiers in the previous Examples, the inventors believed that similar compounds could be used in a similar mechanism to achieve similar advantages.

[0249] The inventors prepared two samples with a ZrB.sub.2 additive to confirm this theory. These samples include: H8838B comprising Fe-31C-3. 7C-0. 3ZrB.sub.2, poured at 1420? C.; and H8843B comprising Fe-30C-4.2C-0.2ZrB.sub.2, poured at 1422? C.

[0250] The optical microstructures of these samples are shown in FIGS. 24(a) and 25(a) and, as shown in FIGS. 24(b) and 25(b), the samples had a lower variance of carbide width, with a median width of about 20-30 ?m and 30-40 ?m respectively. Accordingly, these samples demonstrate a similar effect as the TiB.sub.2 addition in reducing the average primary carbide size, and promoting fine, equiaxed-shaped primary carbides.

Titanium-Based Carbide Modifiers

[0251] The results of Examples 1 and 2 were further validation with samples: H88862A comprising Fe-30C-4.5C-0.2TiB.sub.2, poured at 1420? C.; and H88862B comprising Fe-30C-4.2C-0.4TiB.sub.2, poured at 1420? C.

[0252] The optical microstructures of these samples are shown in FIGS. 26(a) and 26(b), showing a lower variance of fine equiaxed-shaped primary carbides (i.e. limited/no needle-shaped primary carbides). The width of the largest primary carbides and average width of the 20 largest primary carbides were measured, as in Table 3 below.

TABLE-US-00003 TABLE 3 Samples with TiB.sub.2 carbide modifier addition H8862A: H8862B: Fe30Cr4.5C0.2TiB.sub.2 Fe30Cr4.2C0.4TiB.sub.2 AVG AVG Width of width of Width of width of largest top 20 largest top 20 primary primary primary primary carbide carbides carbide carbides Sample (?m) (?m) Sample (?m) (?m) H8862A-1 92 68 H8862B-1 110 68 H8862A-2 91 69 H8862B-2 107 74 H8862A-3 104 68 H8862B-3 127 74 AVERAGE 96 68 AVERAGE 115 72

Conclusion

[0253] Addition of ZrB.sub.2 was found to modify and refine the primary carbides in the hypereutectic HCCI in a manner similar to TiB.sub.2. Accordingly, due to chemical similarities during casting processes, further alternative carbide modifiers may include borides of at least hafnium, vanadium, niobium and tantalum.

[0254] Furthermore, addition of carbide modifiers in the form of a master alloy was found to provide similar advantages as the direct addition of TiB.sub.2 particles or adding Ti and N-containing metals (to form TiN in-situ).

[0255] The mechanics of ZrB.sub.2 as a carbide modifier is believed to be similar to that of TiB.sub.2 demonstrated in the previous Examples. Hence, further carbide modifiers may be the separate addition of: 1) a transition metal source selected from Ti, Zr, Hf, V, Nb and Ta; and 2) a nitrogen source, e.g. in a nitrided steel, such that the sources melt and form a nitride of the transition metal in-situ to function as the carbide modifier.

[0256] Many modifications may be made to the embodiments of the invention described in relation to the Figures without departing from the spirit and scope of the invention.

[0257] In particular, it is noted that the invention is not confined to the modifier of TiB.sub.2, La.sub.2O.sub.3/B mix, ZrB.sub.2 or the modifiers described in relation to the Figures. The invention further extends to, at least, carbide modifiers including borides of Hf, V, Nb and Ta, or in-situ formed nitrides of Ti, Zr, Hf, V, Nb and Ta.

[0258] In the claims which follow and in the preceding description of the invention, except where the context requires otherwise due to express language or necessary implication, the word comprise or variations such as comprises or comprising is used in an inclusive sense, i.e. to specify the presence of the stated features but not to preclude the presence or addition of further features in various embodiments of the invention.