Microstructure and properties of high chromium cast irons: effect of ...

Microstructure and properties of high chromium cast irons: effect of heat treatments and alloying additions E. Karantzalis, A. Lekatou* and H. Mavros Different combinations of critical and subcritical heat treatments variously modify the initial as cast microstructure of high chromium white cast irons leading to secondary carbide precipitation of different extent and nature. Destabilisation (critical heat treatment) of austenite at 970uC for 2?5 h followed by annealing (subcritical heat treatment) at 600uC for 13 h results in massive precipitation of M23C6 carbide particles along with spheroidised M7C3. The reversed order of heat treatments leads to extensive precipitation of M7C3 secondary carbide particles. Mo has a favouring effect on the hardness of the microstructures containing pearlite by limiting pearlite formation. The gradual increase in the alloying additions, C and Cr, increases the hardness of the materials at the different treatment states by inducing carbide precipitation. The increase in the Si content leads to the opposite effect by favouring pearlite formation. Keywords: High chromium white irons, Microstructure, Carbide precipitation, Subcritical heat treatment, Austenite destabilisation

Introduction High chromium white cast irons are extensively used in applications where high friction and wear resistance are demanded, as in ore and mining industry, cement manufacturing and heavy slurry pumping applications. The attractive behaviour they show, especially as far as their wear resistance is concerned, is mainly attributed to their microstructure and the involved phases both in the as cast and after treatment conditions. Tabrett et al.,1 in their extensive review, address all the different parameters that can affect the final microstructure of this family of materials. Their initial cast structure mainly consists of austenitic dendritic matrix along with a eutectic mixture of austenite and carbide M7C3. This initial morphology can be notably modified by the utilisation of various critical and subcritical heat treatments, the purpose of which is the precipitation of secondary carbides and the transformation of the destabilised austenite into more desirable morphologies, such as martensite. The critical heat treatments are usually conducted at 920–1060uC for 1–6 h.1–10 The subcritical heat treatments usually follow the critical heat treatment and are carried out at 200– 600uC for 2–6 h.1–10 Various research efforts have been focused on the formation, morphology and characteristics of the secondary carbides formed during the heat treatments. They have shown that the crystallography, stoichiometry, orientation and extent of formation of these Department of Materials Science and Engineering, University of Ioannina, Ioannina 45100, Greece *Corresponding author, email [email protected]

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ß 2009 W. S. Maney & Son Ltd. Received 2 January 2008; accepted 24 March 2009 DOI 10.1179/174313309X436637

carbides are issues associated with compositional characteristics of the initial material and process parameters during heat treatment.11–21 Despite the nominal high Cr content of these materials, the majority of this is engaged in the carbide structure leaving a Cr deficient matrix with low hardenability.22 Therefore, it is commonly accepted that, in order to increase their hardenability, further alloying elements should be added, especially in cases of large tool dimensions.23 The most common alloying additions are Mo, Ni, Mn and Cu.1–5,23,24 Generally, these additions prevent or enhance the pearlite formation, prevent or enhance the martensite transformation, prevent or enhance the secondary carbide precipitation and, hence, can significantly control the final microstructure and properties.25–39 The present research work is an initial effort to clarify the effect of different heat treatments and alloying additions on the microstuctural characteristics of high chromium white irons and their properties.

Experimental procedure The alloys were prepared by induction melting at 1440– 1460uC and casting in bentonite sand moulds. Hardness was measured in HRC scale with a portable Equotip hardness tester. All castings were subjected to annealing at 600uC for 13 h. A destabilisation treatment at 970uC for 2?5 h followed. Table 1 summarises the compositions of the different alloys. Samples with composition of 2?35%C, 18?23%Cr and 0?58%Mo were subjected to various heat treatments, as shown in Table 2. Metallographic inspection was conducted at a Hund 600 optical microscope. A 10 wt-% ammonium

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Microstructure and properties of high chromium cast irons

2 X-ray diffraction analysis of as cast material (A: austenite; C: carbides; M: martensite) 1 Microstructure of as cast material: austenitic dendrites, eutectic carbides M7C3 and martensitic layers are discerned

persulphate aqueous solution was utilised for chemical etching. X-ray diffraction analysis was carried out using a Bruker D8 advance X-ray diffractometer with a Cu Ka lamp. Scanning electron microscopy examination was conducted by both a Zeiss Supra 35 VP SEM equipped with a Roentec Quantax (Bruker AXs) EDS system and a JEOL 5600 SEM.

Results and discussion Effect of heat treatments As cast alloy

The microstructure of the as cast material is shown in Fig. 1. The alloy consists of large austenitic dendrites along with a eutectic mixture of carbides/austenite. This microstructure has been also observed in previous experimental efforts.1–10 The identified phases are also in agreement with the predicted ones by the Fe rich corner of the Fe–Cr–C phase diagram.1–4 According to Tabrett et al.,1 solidification of such alloys commences with the formation of austenitic dendrites. Then, the eutectic reaction of the lastly solidified liquid follows, which results in the formation of a eutectic mixture of M7C3/austenite. The austenitic phase persists in a

metastable form even at ambient temperatures because of its high content in alloying elements and carbon, the presence of which lowers the Ms temperature below zero. A closer examination of Fig. 1 reveals the existence of a dark martensitic layer surrounding the eutectic carbides. The presence of martensite is explained by the following considerations: the eutectic carbide formation has led to a depletion of the surrounding matrix by carbide forming elements, such as Cr. As a result, the Ms temperature of the austenitic matrix in the vicinity of the carbides increases. Thus, martensite has a worth mentioning presence in the microstructure of the as cast alloy, at room temperature. X-ray diffraction examination (Fig. 2) verifies the limited, yet not negligible, presence of martensite. The major morphological features, in a higher magnification, are shown in Fig. 3a. The large eutectic Table 2 Heat treatment combinations Heat treatment

Conditions

A B

Heating Heating Heating Heating Heating Heating

C D

at at at at at at

970uC 970uC 600uC 600uC 600uC 970uC

for for for for for for

2.5 h, air cooling 2.5 h, air cooling 13 h, air cooling 13 h, air cooling 13 h, air cooling 2.5 h, air cooling

Table 1 Composition of alloys examined Alloy composition, wt-%

Effect of C

Effect of Mo

Effect of Cr

Effect of Si

a/a

C

Si

Mn

Cr

Mo

Ni

Cu

P

1 2 3 4 5 1 2 3 4 1 2 3 4 1 2 3 4

2.67 2.41 2.78 2.45 2.35 2.53 2.53 2.53 2.53 2.56 2.35 2.74 2.23 2.20 2.20 2.20 2.20

0.65 0.65 0.65 0.65 0.65 0.55 0.55 0.55 0.55 0.48 0.48 0.48 0.48 1.60 1.10 0.90 1.20

0.75 0.75 0.75 0.75 0.75 0.68 0.68 0.68 0.68 0.63 0.63 0.63 0.63 0.75 0.75 0.75 0.75

18.10 18.10 18.10 18.10 18.10 18.50 18.50 18.50 18.50 21.29 19.29 23.00 18.09 18.00 18.00 18.00 18.00

0.20 0.20 0.20 0.20 0.20 0.56 0.91 0.40 0.32 0.18 0.18 0.18 0.18 0.21 0.21 0.21 0.21

0.37 0.37 0.37 0.37 0.37 0.33 0.33 0.33 0.33 0.24 0.24 0.24 0.24 0.31 0.31 0.31 0.31

0.70 0.70 0.70 0.70 0.70 0.48 0.48 0.48 0.48 0.48 0.48 0.48 0.48 0.59 0.59 0.59 0.59

0.036 0.036 0.036 0.036 0.036 0.030 0.030 0.030 0.030 0.038 0.038 0.038 0.038 0.040 0.040 0.040 0.040


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3 a microstructure of as cast material (austenitic dendrites and M7C3 plates are recognised) and b EDX spectrum of primary M7C3 eutectic carbide phase

carbide plates (mostly acicular) are surrounded by the austenitic matrix. Energy dispersive X-ray analysis of the primary eutectic carbide phase revealed stoichiometries of the M7C3 type (Fig. 3b). It should be noticed that no secondary carbide formation, as in other works,15 is observed. Heat treatment A (HTA)

Heating of the materials at temperatures as high as 970uC HTA has caused drastic changes in the microstructure. This microstructure is presented in Fig. 4, where large martensitic islands of inner needle-like morphology can be observed in a eutectic carbide/austenite matrix. This needle-like morphology is characteristic of the martensitic transformation. The observed microstructural features have also been identified by previous studies.6–10 The mechanism of martenstic transformation

5 a alloy microstructure after HTA (extensive precipitation of secondary carbides within martensitic grains is observed), b details of circled area in Fig. 6 (carbide precipitates within martensitic grains) and c EDX spectrum of secondary M23C6 carbide phase

for these specific alloys is well established.1,5–12 Apart from these phases, some pearlitic morphologies are also observed, attributed to locally lower cooling rates. Further examination with SEM (Fig. 5a and b) reveals an extensive precipitation of secondary carbide particles within the martensitic matrix. These secondary carbides have a strictly defined cuboid morphology, while their precipitation seems to follow a specific gritlike pattern. Energy dispersive X-ray analysis revealed that the secondary carbide particles have stoichiometries close to M23C6 (Fig. 5c). There is no preferential association of the secondary carbide precipitation with the eutectic carbide phase, while they are almost entirely located within the martensitic grains. These observations are in compatibility with other works,12–21,40,41 where various mechanisms of precipitation are proposed. Especially in the case of the research effort of Powell and Bee,15 the morphology, shape and preferential orientation of the secondary carbide phase are almost

4 Microstructure of alloy after HTA consisting of martensite, eutectic M7C3 and pearlitic morphologies

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6 Alloy microstructure after HTB. ferritic grains, eutectic carbides and pearlitic structures are shown

identical with those observed in the present study. These researchers explained the thermodynamically unfavourable formation of M23C6, in terms of crystallographic matching between the lattice parameters of the carbide and the matrix cubic crystal structure. Such a matching lowers the essential for nucleation activation energy and, thus, reduces the involved surface energy. Heat treatment B (HTB)

Figure 6 illustrates the microstructure of the annealed specimen after HTB. It is seen that annealing has caused significant modifications not only to the microstructure after HTA but also to the microstructure of the as cast material. The microstructure consists of large ferritic grains, eutectic carbide plates and pearlitic structures. A closer examination of the ferritic grains (Fig. 7a) reveals a massive precipitation of secondary carbides, which have been formed either during HTA or during HTB. As shown in Fig. 7b, the secondary carbide particles can be classified into three categories, according to their size: coarse carbide particles, medium carbide particles and fine carbide particles. This specific morphology significantly resembles to that observed by Powell and Bee,15 after prolonged isothermal treatment at the destabilisation temperature (1273 K). Energy dispersive X-ray analysis of the coarse carbide particles suggests stoichiometries of the M7C3 type (Fig. 7c). This is attributed to the fact that, during this heat treatment, some of the M23C6 secondary particles formed during annealing, have been transformed into the thermodynamically more stable M7C3 type. The EDX analysis of the medium size carbide particles revealed a stoichiometry of the M23C6 type. The technique could not provide conclusive results for the fine size carbide particles due to their extremely small size. The presence of other carbide stoichiometries, such as M3C and M6C, is also possible.1,8,12–20 Heat treatment C (HTC)

The prolonged period of annealing in HTC also dramatically alters the microstructure, as observed in Fig. 8. The initial as cast microstructure has been extensively transformed into other morphological features. Three morphologies predominate: pearlitic morphologies, ferritic grains and white colonies of either primary or secondary carbide particles.

7 a alloy microstructure after HTB revealing extensive precipitation of secondary carbide particles within ferritic grains, b alloy microstructure after HTB demonstrating coarse, medium and ﬁne size secondary carbide precipitation and c EDX spectrum of coarse M7C3 secondary carbide phase

Further examination with XRD (Fig. 9) identified the ferritic phase but did not clarify the stoichiometry of the carbide phases, as it is difficult to distinguish between phases, such as M23C6, M6C, M3C, Fe3C, etc., due to extensive peak overlapping. It should also be mentioned that the possibility of martensite and/or austenite presence must not be excluded for the same identification weaknesses. Scanning electron microscopy examination revealed the absence of secondary carbide precipitation of M23C6 or other stoichiometry (Fig. 10). The large white grains are of ferritic nature. Energy dispersive X-ray analysis revealed carbide phases with M/C ratio greater than 7 : 3. Based on these observations, it could be claimed that the primary carbides have been subjected to a combined transformation of spheroidisation and partial degradation to other carbide forms. The overall thermal energy during this heat treatment might not have been sufficient to cause


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11 Alloy microstructure after HTD (three types of microconstituents are discerned: pearlitic colonies, martensitic grains and eutectic carbide acicular plates)

8 Alloy microstructure after HTC (three types of microconstituents are distinguished: pearlitic colonies, ferritic grains and white colonies of either primary or secondary carbide particles)

the formation of Cr rich carbides and pearlitic structures. In addition, they reported martensite formation at intermediate intervals (5, 10 and 15 h) to a different extent. Possible reasons for such differences could be the different alloying element content as well as the different Cr/C ratio. Heat treatment D (HTD)

The final combination of treatments HTD has the most notable effect on the microstructure compared to the other heat treatments. Figure 11 illustrates the microstructure of the alloy after HTD. Three major microconstituents can be distinguished: eutectic carbide acicular plates, a dark phase of acicular morphology most likely being martensite and pearlitic structures. Indeed, XRD spectrum (Fig. 12) revealed the existence of martensite, carbides, retained austenite and ferrite. Further examination with SEM (Fig. 13a) revealed the formation of equiaxed secondary carbide particles of 200 nm mean diameter. It can be seen that the precipitation of these carbide particles has mostly initiated, progressed and integrated within the martensitic matrix. Furthermore, some precipitation has occurred at the periphery of the primary carbide particles. X-ray diffraction analysis was proved inadequate for conclusive results on the secondary carbide stoichiometry. Energy dispersive X-ray analysis revealed stoichiometries of the M7C3 type (Fig. 13b). The formation of martensite during cooling, as aforementioned, is a result

9 X-ray diffraction spectrum of HTC sample (C: carbides; F: ferrite)

10 Alloy microstructure after HTC: large ferritic grains and spheroidised/partially degraded carbides can be observed

extensive phase transformation to other types of carbides but adequate enough to cause reduction of the surface energy by spheroidising and partially transforming the outer carbide layers. These observations constitute a different approach than the approach of Wang et al.,12 who claimed that prolonged annealing (.16 h) at almost similar temperatures (560uC) leads to

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12 X-ray diffraction spectrum of HTD sample (C: carbides; F: ferrite; M: martensite; A: austenite)

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into pearlite, enlargement/spheroidisation of the carbide particles and formation of new secondary carbide particles. The sequence destabilisation treatment, annealing treatment (HTB), is almost entirely adopted in the industrial practice. The prolonged annealing heat treatment (HTC) causes transformation of austenite mainly into ferrite/pearlite. The matrix transformation is accompanied by significant spheroidisation of the primary carbide particles along with the degradation of their outer layers. Secondary carbide precipitation does not occur. The sequence annealing treatment, destabilisation treatment (HTD), causes partial recovery of the initial eutectic carbide morphology, matrix transformation to martensite and significant precipitation of M7C3 secondary carbides.

Effect of alloying additions (C, Mo, Cr and Si)

13 a alloy microstructure after HTD (extensive precipitation of secondary carbide particles within martensitic grains and at periphery of primary carbide particles can be observed) and b EDX spectrum of M7C3 secondary carbide phase

of the Ms temperature increase due to alloying element depletion. The latter was caused by the significant secondary carbide precipitation. Here, it should be noted that several investigations11–16 have reported that the destabilisation isothermal heat treatment mainly leads to extensive precipitation of M23C6 secondary particles. The absence of such precipitation in this effort could be attributed to the fact that the material had been previously subjected to HTC, which, by itself, significantly changes the initial as cast structure. The large spheroidised M7C3 particles observed after HTC have recovered their original acicular shape, whereas they have retained the M7C3 stoichiometry. The outer layers of the original spheroid particles have been dissociated from their cores and transformed by diffusion processes, to nodular secondary carbide particles with a stoichiometry close to M7C3. Transformation to M23C6 has not taken place, since the Cr depleted martensite matrix is very difficult to lead to M23C6 secondary carbide precipitation and prefers the formation of the thermodynamic more stable, Cr poor, M7C3 secondary carbide. Conclusively, a comparison between the microstructures attained by the different heat treatment combinations adopted in this effort could bring forward some interesting morphological issues. The initial as cast microstructure, consisting of metastable austenite and a eutectic mixture of carbide/austenite, transforms to different morphologies depending on the sequence of applied heat treatments. The destabilisation HTA leads to transformation of austenite into martensite accompanied by the extensive precipitation of M23C6 secondary particles. Following HTA, the annealing heat treatment causes significant degradation of martensite

Figures 14 illustrate the effect of different C, Mo, Cr and Si concentrations on the hardness of the alloys in the as cast state and after HTC and HTD respectively. In general, increasing C content leads to an increase in the material hardness irrespective of the state of the alloy. This increase could be mostly attributed to the extent of the primary carbide formed. Other possible synergetic reasons for the hardness increase with the C content could be the increase in the extent of secondary carbide formation, in the strength of the formed martensite and in the strength of the matrix phases. Figure 14a shows that heat treatment D has caused a notable increase in hardness as compared with the as cast state and the HTC. This increase is attributed both to the extensive precipitation of secondary carbide particles and the formation of martensite. Figure 15 presents the microstructure of the sample of the alloy with 2?78%C after HTD. This microstructure does not seem to significantly differ from the morphologies presented so far. It consists of primary M7C3 eutectic carbide acicular plates and martensite. Figure 14b presents the effect of the different Mo additions on the hardness of the alloys in the as cast state, after HTC and HTD. As seen in this figure, the effect of Mo on the hardness appears to be insignificant especially with respect to the as cast and HTD states. Mo has a modestly favouring effect on the alloy hardness after HTC. Mo has been reported to improve hardenability and restrict the pearlite formation.26,27 Heat treatment C has led to microstructures rich in pearlite (Fig. 8). Therefore, it is possible that a decrease in the pearlite content due to Mo is responsible for the increase in hardness observed for the HTC material. Moreover, the hardness of the as cast alloy with 0?91%Mo is higher than that of the lower Mo as cast alloys. This observation is compatible with previous works, where the Mo effect becomes significant at concentrations higher than 1%, even though this statement has not been thoroughly explained.1 Morphologically, the microstructure of these alloys after HTD does alter significantly of what has been presented so far. The XRD spectra of the 0?91%Mo alloy is shown in Fig. 16, and carbides, martensite and ferrite are the identified phases. Table 3 presents the effect of Cr content on the hardness of the alloys as cast and after HTC and HTD. It should be noted that for this specific group of materials, the C content was also changed due to the manufacturing process.


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14 Effect of a C concentration, b Mo concentration, c Cr concentration and d Si concentration on hardness of alloys

Figure 14c presents the Cr effect on the hardness of the alloys. Clearly, the hardness increases with increasing the Cr content irrespective of the alloy heat treatment. Cr is a well established carbide former and ferrite strengthener by solid solution strengthening. Therefore, the strengthening effect of Cr on the white iron can be attributed to the greater extent of primary carbide particles during solidification, the higher amount of secondary carbide particles after HTD and solid solution strengthening of ferritic phases. A clear difference in the microstructure of the alloy with 23?00 wt-%Cr is shown in Fig. 17, especially regarding the extent of the eutectic carbide phase and martensite. The higher the Cr content, the more restricted the martensitic phase and the more extended the primary carbide phase. As with C, by increasing Cr content the eutectic point of the Fe–Cr–C system lowers down. Thus, the percentage of the primary carbide phase increases at the expense of the martensitic phase. The XRD spectrum of the 23?00%Cr sample (Fig. 18) after HTD revealed the presence of carbides, martensite, ferrite and austenite. The presence of austenite is most likely associated with the high Cr content. As a carbide former, Cr depletes austenite of C. Thus, the Ms temperature increases.

15 Alloy microstructure with 2?78%C after HTD: eutectic carbide plates and martensitic matrix can be observed; pearlitic structures may occasionally be present

Table 3 Hardness modiﬁcation as function of Cr (and partially C) content Hardness, HRC Sample Cr–C contents, wt-% Cr/C As cast HTC HTD 1 2 3 4

16 X-ray diffraction spectrum of 0?91%Mo sample after HTD (C: carbide M7C3; M: martensite; F: ferrite)

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18.09–2.23 19.29–2.35 21.29–2.35 23.00–2.74

8.11 8.21 8.31 8.39

48.3 49.1 54.6 55.6

37.9 39.1 42.7 44.1

64.5 65.6 66.1 67.1

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20 X-ray diffraction analysis of 1?6%Si sample (C: carbides; M: martensite; F: ferrite) 17 Alloy microstructure with 23?00%Cr after HTD: eutectic carbides and martensitic matrix can be observed

Figure 19 shows the microstructures of the sample with 1?6 wt-%Si, which it contains a significant amount of pearlitic colonies. X-ray diffraction analysis of the 1?6%Si alloy (Fig. 20) revealed the presence of martensite, carbide (M7C3) and ferrite.

Conclusions

18 X-ray diffraction of 23?00%Cr sample after HTD (C: carbides; M: martensite; F: ferrite; A: austenite)

19 Alloy microstructure with 1?6%Si after HTD, consisting of eutectic carbide plates, martensite, ferrite and pearlite

Figure 14d illustrates the effect of different Si additions on the hardness of the alloys. Si adversely influences the hardness of the alloys irrespective of the heat treatment. This could be owing to the tendency of Si to stabilise the softer pearlitic structures.37,38 Moreover, this could be the most possible reason why the hardness lowering effect is stronger in the case of HTC. Heat treatment C leads to extensive pearlite formation in comparison with the other heat treatments.

1. The initial as cast structure of a high chromium white iron consisted of primary austenitic dendrites and a eutectic mixture of M7C3 plates/austenite. Cr depletion of the matrix area in the close vicinity of the carbide particles has led to their engulfment by a thin martensite layer. 2. The destabilisation heat treatment (HTA) drastically altered the microstructure, leading to extensive cubic carbide particle precipitation, of M23C6 type, within the matrix. The latter had been transformed into martensite. 3. The subsequent subcritical heat treatment (the combination of destabilisation–subcritical compose HTB) led to the transformation of the martensitic matrix to ferrite/carbide in dispersion. The overall secondary carbide distribution was significantly altered and consisted of coarse M7C3 particles, medium M23C6 particles and fine unidentified carbide particles. 4. Direct subcritical heat treatment (HTC) of the as cast material caused the formation of pearlitic structures along with the spheroidising/outer surface degrading of the primary eutectic M7C3 particles. No secondary carbide precipitation was observed. 5. The subsequent destabilisation process (the reverse combination of subcritical destabilisation compose HTD) led to partial regaining of the initial eutectic carbide morphology possibly by dissociation of the outer unstable layers. The latter had been transformed to nodular secondary M7C3 particles that had precipitated in the close vicinity of the primary carbide particles. Pearlite was transformed into martensite. 6. Mo additions lower than 0?91 wt-% do not significantly affect the hardness of the alloy in the as cast state and after HTD. Regarding the HTC state that leads to pearlite formation, Mo seems to increase the hardness by limiting pearlite formation. C and Cr additions increase the hardness of the alloy mostly by inducing primary and secondary carbide precipitation, at the expense of martensite. Si additions decrease the hardness of the alloy by favouring pearlite formation.


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Acknowledgement The authors would like to acknowledge the Greek Modern Castings (GMC) SA foundry for their kind assistance in the preparation of the examined alloy.

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