Mechanical Properties
Effect of Si Content on the Microstructure and Wear Resistance of High Chromium Cast Iron
2018 Volume 58 Issue 8 Pages 1532-1537
Details
2018 Volume 58 Issue 8 Pages 1532-1537
Effect of Si content on the wear resistance of high chromium cast iron was investigated. The result showed that the eutectic carbides were greatly refined and increased with the increase of Si addition from 0.5 to 1.5 wt.%. Meanwhile, the transformation of austenitic matrix to pearlite was observed. After destabilized treatment at 950°C for 3 h, denser secondary carbides precipitated from the matrix are observed in the alloy containing 1.5wt.%Si. Such microstructure changes leaded to the increase of hardness and improvement of abrasive wear resistance. An increase in load from 20 N to 100 N, the extent of carbide fragmentation was much greater and extended to a deeper distance. Severer fracture of carbides was observed in the alloy containing 0.5wt.%Si. The higher resistance with addition of Si was attributed to the denser secondary carbides, which strengthened the matrix between the eutectic carbides and in turn increased the mechanical support to the carbides against cracking.
The high chromium cast iron has good wear resistance and is widely used in slurry pumps and hot strip rolls due to its good wear resistance.1,2,3) The excellent wear resistance of the high chromium cast iron is related to the high volume fraction of the hard M7C3 carbides. However, the high volume of M7C3 carbides inevitably decreases of fracture toughness due to its brittleness.4,5)
The wear resistance of high chromium cast iron is affected by both the matrix and the carbide characteristics (type, morphology and size). Fulcher et al.,6) suggested that the matrix had a synergetic protection effect with the eutectic carbides during wear process. The wear resistance of the cast iron depended on the level of protection which the carbides provided to the matrix each other. Generally, the matrix can be strengthened by applying destabilization treatment to get a harder martensitic matrix with secondary carbides distributing in it.7)
Titanium and niobium were used to improve the wear resistance by forming the harder carbides.8,9,10) Different from the titanium and niobium, silicon was rejected from forming carbides and mostly dissolved in the matrix. In the medium and high carbon high strength steel by applying quenching-portioning -tempering treatment (QPT), the addition of Si with high content (1.5–2.0 wt.%) was used to suppress the precipitation of cementites during portioning process and enhance the carbon partitioning from martensite to austenite to increase the stability of austenite.11,12) In the high chromium cast iron, silicon was reported to refine the eutectic carbides colonies and destabilize the austenite, resulting in the transformation of austenite to pearlite during solidification.13,14) Lai et al.15) found that the addition of Si can promote the precipitation of secondary carbides during destabilization process. Therefore, it can be postulated that the wear resistance of high chromium cast iron should be improved by the addition of silicon without at the cost of decreasing the toughness. The effect of Si on the abrasion wear resistance has not yet been studied. In this paper, the effect of Si on the microstructure and abrasion wear resistance of high chromium cast iron was investigated.
The actual chemical composition of the high chromium cast irons in the present work was listed in Table 1. The alloys investigated were melted in a medium-frequency induction furnace and cast at into a cylindrical metal mould with dimension of Φ200×300 mm. Specimens for microstructural observation and abrasive wear tests were cut from the cast bars using wire-electrode cutting machines. Figure 1 shows a schematic drawing of retrieving specimens from the cast bars for microstructural analysis and abrasive wear test.
| alloy | C | Si | Mn | Cr | Mo |
|---|---|---|---|---|---|
| A | 2.81 | 0.50 | 0.61 | 18.1 | 0.85 |
| B | 2.82 | 1.50 | 0.62 | 18.2 | 0.83 |
A schematic drawing of retrieving specimens from the cast bars.
All specimens were destabilized at 950°C for 3 h, followed by quenching in the air at room temperature and then tempered at 200°C for 6 h to get a tempered martensitic matrix and released stress. Specimens for microstructural observation were mechanically polished and then etched by a 3% nital solution. Scanning electron microscope (SEM, Sirion 200) was used to characterize the microstructure and the worn surface after abrasive wear test. Chemical composition of different phases was determined using the electron probe micro-analyzer (EPMA, JMA-8230). The microhardness tests were carried out on a microhardness tester (BUEHLER 5104) by a 20 g load for 15 s and each data given represent the average of at least five measurements.
A schematic drawing of the abrasion wear testing machine was shown in Fig. 2. The three-body abrasive wear test was performed on the MLS-225 rubber wheel abrasive tester according to the ASTM G105-16 standard.16) The slurry was a combination of 1000 mL of distilled water and 1500 g of sand (50–70 mesh silica sand). The specimens with dimension of 55×25×6 mm was mechanically polished before abrasive wear test. The test was carried out at a rotating speed of 240 rpm for 25 min and under loads of 20, 60, 100 N. The weight loss was measured on an analytical balance with a resolution of 0.001 g.
A schematic drawing of the abrasion wear testing machine.
Figure 3 shows the as-cast microstructure of the alloys at positions A and B in Fig. 1, respectively. Both the alloys are hypoeutectic microstructure, consisting of the primary phase (retained austenite/pearlite) and eutectic carbides colonies. The eutectic carbides had a hexagonal crystal structure and exhibited different structure with different orientations. The carbides appeared as chrysanthemum-like colonies in the horizontal direction and rod-like laths in longitudinal direction. For the alloy B, the microstructure presents finer eutectic carbide colonies compared to that of the alloy A. The length of the eutectic carbides is determined by measuring the rod of a selected rod-like carbides volume in the Figs. 3(c) and 3(d). For example, the random carbides colonies are measured to be 243 μm for the alloy A and 123 μm for the alloy B, as shown in Figs. 3(c) and 3(d). The average length of the eutectic carbides measured is decreased from 220 to 110 μm with the increase of Si from 0.5 to 1.5 wt.%. The amount of carbides is increased from 30±2% for the alloy A to 35±2% for the alloy B by using Leica digital images analyzer. Similar result was also obtained by A Beddolla Jacuinde et al.8) Meanwhile, the austenitic matrix was transformed into pearlitic matrix with increase of Si addition. The formation of pearlite favored by addition of Si was reported by Laird II et al.17)
As-cast microstructures of the alloys: (a, c) alloy A at positions A and B, respectively; (b, d) alloy B at positions A and B, respectively.
Figure 4 shows the SEM images of as-cast microstructures of the alloys using the second electron mode (SE), revealing the austenitic matrix in the alloy A and the pearlitic matrix in the alloy B. Further examination of the matrix with higher magnification showed that the pearlitic matrix consisted predominantly of lamellar pearlite, as shown in Figs. 4(c) and 4(d). Since EDS cannot determine the precise concentration of the elements with low atomic numbers, e.g. carbon, the chemical compositions of the matrix with different phases are determined by the electron probe microanalysis technique (EPMA) with higher precision, as listed in Table 2. The pearlitic matrix of the alloy B possesses higher carbon, silicon and chromium content than the retained austenite of the alloy A. The mechanism of the transformation of austenitic matrix to pearlite can be inferred as follows: the increase of Si content increases the solid solubility of carbon in matrix and lowers the stability of austenite at high temperature. As a result, the decomposition of austenitic matrix to pearlite was found.
| alloy | Matrix | C | Si | Mn | Cr | Mo | Fe |
|---|---|---|---|---|---|---|---|
| A | Retained austenite | 3.6±0.1 | 1.5±0.1 | 0.4±0.05 | 12±0.5 | 2.5±0.1 | Bal. |
| B | Pearlite | 11±1.0 | 2.2±0.1 | 0.4±0.05 | 15±0.5 | 2.5±0.1 | Bal. |
The improvement in the wear resistance of high-Cr cast iron can be achieved when the as-cast austenite matrix was transformed to martensitic matrix by destabilization treatment.18) The martensitic matrix with precipitation of numerous secondary carbides is often preferred for many wear applications.17) The microstructures of the alloys exhibiting the matrix and distribution of secondary carbides after destabilizing treatment are shown in Fig. 5. The right images (Figs. 5(b), 5(d)) are the magnification of the red rectangle in the left images (Figs. 5(a), 5(c)) to illustrate the distribution of the secondary carbides more clearly. It can be seen that denser secondary carbides precipitated were observed in the matrix of the alloy B. The microhardness of matrix increased from 860 to 920 HV, which is consistent with the increase of macrohardness from 58.5 to 61. HRC with the increase of Si content, as listed in Table 3. The increase of macrohardness is associated with the higher microhardness of matrix resulting from the precipitation of denser secondary carbides and the increase of ~5% the amount of eutectic carbides. Higher carbon and chromium content in as-cast matrix of alloy B provided more elements concentration source for the precipitation of denser secondary carbides like M7C3 or M23C6 during destabilization treatment.
Microstructures after destabilizing treatment: (a, b) alloy A; (c, d) alloy B.
| alloy | Microhardness of matrix (HV) | Macrohardness (HRC) |
|---|---|---|
| alloy A | 860±10 | 58.5±1 |
| alloy B | 920±10 | 61.2±1 |
The wear test was performed according to the ASTM G105-16 standard.16) Figure 6 shows the abrasive mass loss of the alloys after destabilization treatment with different loads. It can be seen that the abrasive mass loss increases with the increases of the load. The effect is more pronounced at the load range from 60 to 100 N. For all loads, the alloy B presents less abrasive mass loss than the alloy A, indicating the better wear resistance.
The abrasive mass loss as a function of the load.
Figure 7 shows the SEM images of worn surfaces of the alloys. The secondary electrons (Figs. 7(a), 7(c)) and back-scattered electrons (Figs. 7(b), 7(d)) of the SEM technique are used to characterize the worn surfaces of the alloys. According to the JJ Penagos et al.,19) the mechanism of three-body abrasion is a combination of micro-cutting with plastic deformation. For the alloy A, the worn surface shows deeper grooves and more pitting sites, suggesting a higher severity of wear than the alloy B. The back-scattered electron imaging shows that the pitting occurs preferentially at the eutectic carbides, causing the cracking of eutectic carbides. Higher magnification of the worn surface is employed to analyze the depth of the grooves and roughness of the worn surface, as shown in Fig. 8. It can be seen that the worn surface of the alloy A showed deeper and wider grooves than that of the alloy B, exhibiting a rougher surface.
SEM images of worn surfaces with load of 100 N: (a, b) alloy A; (c, d) alloy B.
SEM images of localized grooves on the worn surface: (a) alloy A; (b) alloy B.
Cross sections were taken from the worn surface to investigate the fracture and cracking of the carbides. Figure 9 showed the cross-section of the alloys when tested at 20 N. No obvious differences in the fracture of the carbides were found in both alloys. Only localized subcritical cracking was found in some carbide in both alloys, as arrowed in Fig. 9. The position where the carbides crack is slightly deeper in the alloy A (~6 μm) than that of the alloy B (~4 μm). When tested at higher load of 100 N, the extent of carbide fragmentation was much greater and extended to a deeper distance (11–17 μm), as shown in Fig. 10. However, the level of cracking showed much difference in both alloys. For the alloy A, the carbides are completely cracked and extended to a deeper distance close to the surface, while the most of the cracks of the alloy B are sub-critical. The difference of the cross section micrographs between the loads of 20 N and 100 N is consistent with the result of abrasive mass loss, as shown in Fig. 6.
SEM images of the cross-section of the alloys tested at 20 N: (a) alloy A; (b) alloy B.
SEM images of the cross-section of the alloys tested at 100 N: (a, b) alloy A; (c, d) alloy B.
The higher wear resistance of the alloy B can be explained as below: higher Si content promoted the precipitation of denser secondary carbides, which can strengthen the matrix and increase the hardness. According to the synergetic protection theory of matrix and carbides,6) the role of carbides to resist the abrasive grits relies on the protection that the matrix provide to it. Therefore, the matrix with higher hardness in the alloy B can provide better protection effect to the carbides, resulting in higher wear resistance. On the contrary, the carbides of the alloy A are not well protected, resulting in much severer fracture.
(1) Eutectic carbides are greatly refined and increased. The transformation of austenitic matrix to pearlite is observed with the increase of Si addition from 0.5 to 1.5 wt.%. Denser secondary carbides are observed in the alloy containing 1.5wt.%Si.
(2) The alloy containing 1.5wt.%Si shows better wear resistance. An increase in load from 20 N to 100 N, the extent of carbide fragmentation was much greater and extended to a deeper distance. Fully fracture of carbides is observed in the alloy containing 0.5wt.%Si, while the carbides in the alloy containing 1.5wt.%Si show only sub-critical cracking.
This work is financially supported by the Innovation Foundation of Central South University (No. 2016zzts028).
