Effects of Primarily Solidified Dendrite and Thermal Treatments on the M₂₃C₆ Precipitation Behavior of High-Chromium White Iron

Abstract

The precipitation behavior of M₂₃C₆ carbide during thermal treatment of high-Cr white iron with various fractions of primarily solidified dendrite was studied and reviewed. M₂₃C₆ precipitation in the primarily solidified dendrite occurred preferentially during conventional heat treatment, whereas it occurred scarcely in the eutectic austenite. The reaction between M₇C₃ and austenite caused the dissolution of M₇C₃ into austenite, followed by precipitation of M₂₃C₆ along the periphery of eutectic M₇C₃. Relatively low-temperature thermal treatment (modified heat treatment) led to precipitation of M₂₃C₆ particles in the eutectic austenite, which is presumed to be caused by solubility difference depending on temperature.

1. Introduction

The superior wear resistance of high-Cr white iron enabled the application of this material in mining activities such as ore grinding, coal grinding, or in their transport. The superior wear resistance of these alloys is originated in their matrix microstructure and the existence of various carbides, both in as-cast and heat-treated conditions. As-cast high-Cr white irons may have martensitic or austenitic matrix depending on the post-casting cooling rate and section size. The heavy sections that cool slowly have a martensitic matrix with a small amount of relatively soft-retained austenite due to incomplete transformation to martensite. The maximum hardness of the matrix is achieved by the subsequent heat treatment that transforms the matrix to martensite. However, the pearlite matrix has less hardness and low toughness; thus, pearlite is not desirable. Pearlite formation occurs when alloying is insufficient [1].

The matrix structure and existing carbides have strong effects on the mechanical properties of high-Cr white cast irons. With a process similar to the solidification of hypoeutectic gray irons [2], hypoeutectic high-Cr white irons solidify primarily with the formation of austenite dendrites, followed by a eutectic reaction between austenite and carbide. The carbides in high-Cr white iron are very hard and wear resistant but very brittle [1]. Wear resistance could be improved by increasing the amount of carbide.

The types of existing carbides in high-Cr irons are M₇C₃, M₃C, and M₂₃C₆. Generally, the eutectic reaction generates M₇C₃ in high-Cr white iron. M₇C₃ carbides contribute to improving the wear resistance of the alloy; however, those primarily precipitated from the melting process ahead of eutectic reaction are known to be quite deleterious to impact toughness and should be avoided [1]. Thus, the hypereutectic composition of the alloys is not desirable for engineering applications. In hypoeutectic alloys, increasing carbide fraction decreases toughness. However, the higher fraction of austenite or dendrite may improve toughness but reduce the hardness of the alloy. It is known that the fractions of austenite and M₇C₃ carbide in hypoeutectic high-Cr white iron are sensitive to C and Cr contents [1]. Decreasing the amount of C and Cr increases the austenitic dendrite fraction and decreases the fraction of M₇C₃ along the interdendritic regions.

The release of C and Cr from the saturated austenite and the precipitation of M₂₃C₆ carbide, which is known as “destabilization”, could occur at high temperatures. The degree of destabilization at high temperatures is expected to be related to the fraction of dendrite. Variation in dendrite fraction through modification of chemical compositions also means various fractions of eutectic structure. As reported earlier, the dendritically solidified austenite has a slightly different chemical composition from the eutectically solidified one [3,4]. Thus, the destabilization behavior in dendritic austenite might be different from that in eutectic austenite. M₂₃C₆ precipitation temperature in the alloys that have various fractions of existing phases might be different from each other. Therefore, the purpose of this study was to understand the precipitation behavior of M₂₃C₆ or destabilization of austenite in the hypoeutectic high-Cr white irons with various dendrite fractions. In particular, the modified heat treatment condition (direct aging (DA)), in which the maximum fraction of M₂₃C₆ precipitation occurred for each alloy, was selected on the basis of ThermoCalc prediction. Precipitation behavior of M₂₃C₆ in austenite (with both dendritic and eutectic structures), and the reaction between austenite and M₇C₃ during the modified heat treatment were studied.

2. Experimental Procedure

2.1. Specimen Preparation

As reported earlier by the authors [5], the primarily solidified dendrite fractions in high-Cr white iron were adjusted in the composition range of 2.1~2.9 wt.% C and 24.0~27.0 wt.% Cr, which is in the composition range of ASTM A532 Class 3 Type A. Chemical analysis of the alloys was carried out by optical emission spectrometer (OBLF QSN-750) and carbon–sulfur determinator (ELTRA CS800). The chemical compositions of the specimens are given in

Table 1. Chemical composition of the experimental high-Cr white iron alloys [5].

Designation	Composition (wt.%)
	C	Si	Mn	Cr	Ni	Mo
2124	2.12	0.66	0.65	24.05	0.85	0.86
2127	2.13	0.71	0.67	27.00	0.89	0.86
2427	2.43	0.68	0.70	27.06	0.86	0.85
2827	2.78	0.70	0.70	27.34	0.87	0.85
2927	2.95	0.67	0.71	26.94	0.92	0.82

. Phase prediction was carried out to understand the phase evolution during cooling, with the chemical composition provided in

Table 2. Precipitation behaviors of M₇C₃ and M₂₃C₆ by ThermoCalc calculation [5].

Alloy	M23C6 Precipitation Initiation		M23C6 Peak Precipitation
	T (°C)	M7C3 Fraction (mol.%)	T (°C)	M23C6 Fraction (mol.%)	M7C3 Fraction (mol.%)
2124	1060	19.8	850	17.3	12.2
2127	1180	18.8	890	33.0	2.5
2427	1100	23.4	870	20.1	14.0
2827	1000	29.1	810	9.8	24.6
2927	920	31.9	739	6.0	29.7

through commercial software ThermoCalc (Thermo-Calc 2019b) on the basis of database DB TCFE9 for Steels/Fe-Alloys (V.9.1). The specimens were prepared by induction melting of the master alloy, followed by secondary melting and casting to a rod.

The cast specimens were usually subjected to conventional heat treatment and modified heat treatment. The conventional heat treatment condition included soaking at 1065 °C for 4 h, followed by air cooling and tempering at 500 °C for 4 h, and 250 °C for 4 h, respectively [5]. To understand the precipitation behavior of M₂₃C₆, a modified heat treatment condition for each alloy was selected on the basis of ThermoCalc prediction. The modified heat treatment (DA) temperature for each alloy was chosen as the temperature at which the maximum M₂₃C₆ fraction appeared on ThermoCalc calculation. The soaking time at each DA temperature was decided by the time (4 h) at which the highest hardness value was shown. Rockwell hardness values (WOLPERT / D-6700) of the specimens were measured with 10 points in the as-cast, conventionally heat-treated, and modified heat-treated conditions. The DA treatment included soaking at each temperature for each alloy, followed by air cooling in the absence of the tempering process.

The details of conventional heat treatment and modified heat treatment (direct aging (DA)) are as follows:

• Conventional heat treatment of destabilizing and tempering

  -

  1065 °C for 4 h air cooling (AC) + 500 °C for 4 h (herein: conventional heat treatment);

  -

  1065 °C for 4 h water quench (WQ) + 500 °C for 4 h (herein: DES + WQ + Temp).

• Modified heat treatment (direct aging (DA)): Each alloy was subjected up to 240 min at the temperature of maximum fraction of M₂₃C₆. The DA treatment of each alloy is as follows:

  -

  2124 alloy: 850 °C for 10, 30, 60, 240, 300 min + AC;

  -

  2127 alloy: 890 °C for 10, 30, 60, 240, 300 min + AC;

  -

  2427 alloy: 890 °C for 10, 30, 60, 240, 300 min + AC;

  -

  2827 alloy: 810 °C for 10, 30, 60, 240, 300 min + AC;

  -

  2927 alloy: 800 °C for 10, 30, 60, 240, 300 min + AC.

• DA specimens were exposed under 1 × 10⁻⁵ torr vacuum for 60 min to observe the movement of the M₇C₃ carbide interface. Certain points of DA specimens were engraved to observe at the same position before and after the vacuum exposure. The vacuum exposure was carried out in a vacuum furnace (Jungmin 2010), followed by Ar gas fan quenching.

2.2. Microstructural Observation and Phase Identification

Specimens for optical microscopy (OM: Nikon ECLIPSE MA200) and scanning electron microscopy (SEM: JEOL JSM IT500LV) were prepared metallurgically and etched by swabbing with Vilella’s reagent consisting of 45 mL glycerol, 15 mL nitric acid, and 30 mL hydrochloric acid. The volume fraction of the primarily solidified dendrites in the as-cast specimens was measured during optical microscopy by an Image analyzer (iMT iSolution DT). Specimens for transmission electron microscopy (TEM: JEOL JEM-2100F) were prepared by mechanical polishing down to 60-micrometer thickness, followed by twin jet polishing (Struers TenuPol-5). The solution for twin jet thinning was 10% perchloric acid in methanol. The existing phases in the heat-treated specimens were identified by TEM selected area diffraction pattern (SADP), and energy-dispersive X-ray spectrometer (EDS: Oxford X-MAX AZtec).

3. Results and Discussions

3.1. Prediction and As-Cast Microstructure

As reported earlier [6,7], the M₇C₃/austenite eutectic phase and primarily solidified dendrite coexist in the as-cast condition of high-Cr cast irons. The eutectic carbide in the white iron with more than 20% Cr content is M₇C₃ [1]. Similar to gray irons, the primarily solidified dendrite forms during the solidification of high-Cr white irons. Oh et al. [5] reported various fractions of primarily solidified dendrite in high-Cr white irons, as shown in Table 1. It is known that the primarily solidified dendrite formation during solidification of cast irons is closely related to C equivalent value (Ceq). The equilibrium microstructural evolution during post-casting was able to be predicted through ThermoCalc calculation, as shown in Figure 1. Solidification began with austenite (or delta ferrite) in the alloys with low-Ceq-valued hypoeutectic high-Cr white iron, followed by the eutectic reaction of austenite and M₇C₃, as shown in Figure 1. The solidification of the austenite (or delta ferrite) forms dendrites, which are referred to as primarily solidified dendrites in this study. As displayed in Figure 1, phase transformation and precipitation of new phases occurred during the post-casting stage. It was possible to predict the precipitation behavior and phase transformation during cooling through commercial software ThermoCalc.

In the as-cast microstructure of the alloys [5], the primarily solidified dendrite fraction increased with decreasing Ceq, as predicted in ThermoCalc calculation. The lowest Ceq alloy, 2124 had a high fraction of dendrite, and little dendrite was found in the highest Ceq alloy, 2927.

3.2. Conventional Heat Treatment

Conventional heat treatment of the high-Cr cast iron included soaking at 1065 °C, followed by air cooling and tempering. During soaking, “destabilization”, which allows the solute elements such as C and Cr to be derived from the saturated austenitic matrix to precipitate the secondary carbides, preferentially occurs [5]. M₂₃C₆ carbide precipitation from the saturated austenite during destabilization is expressed as follows in the previous report [4,5].:

γ → γ * + M₂₃C₆

(1)

where γ* is austenite with lower alloy content than that of the original matrix γ [4,8].

The conventionally heat-treated alloys have dendrite features similar to those in as-cast alloys [5], but dendritic boundaries were not distinct compared with as-cast types. Compared with those of the as-cast alloys, the dendrites were filled with many particles in the conventionally heat-treated condition. Little or no particles exist within the dendrites in the as-cast condition [5], as displayed in Figure 2 and Figure 3; however, conventionally heat-treated alloys have a considerable number of precipitated particles in the dendrites, as displayed in Figure 4. The particles are secondary M₂₃C₆ carbide [4]. In the present study, TEM micrographs and selected diffraction pattern analysis identified the particles to be M₂₃C₆, as shown in Figure 5. The diffraction patterns in Figure 5 show that most of the particles had an orientational relationship with the matrix.

The micrographs of dendritic and eutectic austenites in conventionally heat-treated conditions are different from each other. As mentioned earlier, dendritic austenite was filled with M₂₃C₆ particles, while little or no M₂₃C₆ particles existed in the eutectic austenite, as displayed in Figure 6 and Figure 7. This means the nature of austenite with a dendrite structure is different from that with a eutectic structure. As Dupin showed schematically, the alloying elements are believed to contribute to forming M₇C₃ carbide during a eutectic reaction; thus, C and Cr in the austenite enveloping the M₇C₃ carbides are depleted [3,4]. As the austenite in the dendrite formed before the eutectic M₇C₃/austenite reaction, the solute concentration of the austenite in dendrite might be rich compared with that of the eutectic austenite. This means that the primarily solidified dendrites have a high potential to form the precipitates due to saturation of solute elements such as Cr and C. Therefore, easier and more precipitation of M₂₃C₆ occurred during destabilization in the dendritic austenite than in the eutectic austenite.

Except for 2927 alloy (eutectic alloy), precipitation of M₂₃C₆ scarcely occurred in the interdendritic eutectic area. Many studies also explained that the secondary carbides do not nucleate and grow on eutectic carbides but form preferentially within the dendritic matrix [9,10,11]. The present study found a similar result in hypoeutectic alloys, as displayed in Figure 6 and Figure 7. When comparing Figure 6 and Figure 7, the precipitation is strongly related to the existence of primarily solidified dendrites in the hypoeutectic alloys, as reported in earlier studies [5,9,10,11].

However, the dendrite-free and fully eutectic 2927 alloy had a little amount of M₂₃C₆ carbides in the austenite, which formed during the eutectic reaction, as displayed in Figure 7e. This is presumed to be related to the solubility of the alloying elements depending on temperature, which is addressed in direct aging treatment (DA).

3.3. Modified Heat-Treated (DA-Treated) Microstructure

3.3.1. Basic Background of M₂₃C₆ Precipitation

According to the ThermoCalc prediction, the fraction of M₂₃C₆ carbide increases at the cost of M₇C₃ and austenite reduction at high temperatures. The fraction of M₂₃C₆ carbide in each alloy was inversely proportional to that of M₇C₃ carbide, as shown in Figure 1. Previous studies, mentioned above, revealed that M₂₃C₆ carbides precipitated during destabilization; thus, precipitation preferentially occurs in the dendrites [5,9,10,11]. This also occurred in the present study.

Then, a question arises from the ThermoCalc results: How does the fraction of M₂₃C₆ increase with the reduction of M₇C₃ and austenite? This is shown in Figure 1 and Table 2. The precipitation-initiating temperature of M₂₃C₆ carbide during post-cast cooling is listed in Table 2 for each alloy, in comparison with its volume fraction and the temperature of maximum M₂₃C₆ fraction. The phase fraction of M₂₃C₆ carbide in the alloy 2927 (high Ceq and high fraction of M₇C₃) was relatively low, compared with that of the other alloys. The amount of M₂₃C₆ carbide was less in high-Ceq alloys than in low-Ceq alloys (with a high fraction of dendrite). The amount of M₇C₃ reduction in the high dendrite fraction alloys (low-Ceq alloys 2124, 21,217, and 2427) was more than in high-Ceq alloys (2827 and 2927).

Additionally, another question arises: Is M₂₃C₆ formed only due to the transformation [9,10,11] of M₇C₃ to M₂₃C₆ occurring at high temperatures? If so, alloys with a high fraction of M₇C₃ have more possibility of transition than alloys with a low fraction of M₇C₃. Hence, the reduction in M₇C₃ in the alloys with a high fraction of M₇C₃ would be more than that in the alloys with a low fraction of M₇C₃ during post-casting cooling or under equilibrium conditions. However, the alloys with the low fraction of M₇C₃ (high fraction of dendrite) showed a relatively high fraction (including destabilized carbide) of M₂₃C₆, as listed in Table 2. Similar to the findings Pearce reported earlier [11], the result showed that the formation of M₂₃C₆ may not result only from the transformation from M₇C₃ to M₂₃C₆ but also from the precipitation from the adjacent austenite.

The fractions of M₇C₃ and M₂₃C₆ did not always have a linear relationship with Ceq; for instance, the M₇C₃ fraction of 2127 was the lowest among the alloys, whereas that of M₂₃C₆ was highest. This might be related to the initiation of solidification with delta ferrite in the alloy. Delta ferrite in the alloy might contain a high concentration of Cr and C; thus, 18.1% of M₇C₃ fraction at the fully solidified temperature (1285 °C) increased to 19.5% at the cost of delta ferrite reduction (from 9.9% at 1284 °C to 0 at 1205 °C). Below this temperature, M₂₃C₆ precipitation increased due to the destabilization that allowed Cr and C to arise from the austenite matrix, as well as the reduction in a small portion of M₇C₃. The fraction of M₇C₃ decreased to 2.4% at 900 °C, at which maximum precipitation of M₂₃C₆ occurred. In particular, the fraction of M₂₃C₆ in 2127 alloy was the highest, and that of M₇C₃ was the lowest among the experimental alloys. Solidification of 2127 alloy begins with delta ferrite before austenite solidifies. The relatively small fraction of M₇C₃ in 2127 alloy might be related to the delta ferrite. Additionally, it is possible to suppose that delta ferrite might have a high concentration of Cr and C. The consumption of Cr and C at the early stage of solidification may cause deficiency of the elements in the remaining liquid for sufficient eutectic reaction. Thus, the alloy (2127) may have less amount of eutectic reaction compared with the other alloys. This would be indirect evidence indicating that delta ferrite has a relatively high content of Cr. Unfortunately, it was impossible to measure the Cr concentration in delta ferrite due to the rapid transformation of delta ferrite to austenite and finally to martensite. However, based on the ThermoCalc calculation and microstructural observations, the above assumption of early consumption of Cr and C during delta ferrite formation might have caused a deficiency in solute elements in the remaining liquid; thus, a relatively small amount of M₇C₃ and eutectic reaction occurred during the final freezing. In fact, the fraction of M₇C₃ in 2127 was lower than that of the other alloys. The difference in M₂₃C₆ precipitation among the alloys can be found in Figure 8 and Figure 9 (DES + Temp of both AC and WQ). As listed in Table 2, the fraction of M₂₃C₆ was the highest in 2127 alloy, while it was medium in 2124 and 2427, and low in 2827 and 2927 because of having little or no dendrites when destabilization occurred.

3.3.2. Effect of Direct Aging (DA) on the M₂₃C₆ Precipitation

On the basis of ThermoCalc calculation, direct aging of each alloy in the absence of destabilizing treatment or solutionizing was conducted at the temperature of the maximum fraction of M₂₃C₆, as shown in Table 2. Unlike conventionally heat-treated alloys, DA-treated alloys had M₂₃C₆ precipitates in both dendritic and eutectic austenites of M₇C_3, as displayed in Figure 10. This is significantly different from the conventionally heat-treated alloys, which are displayed in Figure 8 and Figure 9, where M₂₃C₆ precipitation occurred in the dendrite region of austenite but not in eutectic austenite. As mentioned above, M₂₃C₆ precipitation during the conventional heat treatment is mainly caused by the destabilization of the saturated austenite, as expressed in Equation (2). This also leads to a difference in composition (or degree of saturation) between the dendrite austenite and that in the eutectic region. The primarily solidified austenite may have a saturation of solute elements such as Cr and C; however, the austenite in the eutectic region may have fewer solute contents due to Cr-rich M₇C₃ carbide formation during the eutectic reaction.

It is known that the transition or replacement of carbides from metastable M₇C₃ to stable M₂₃C₆ occurs during destabilization heat treatment [11,12]. It results in M₂₃C₆ shells surrounding the eutectic M₇C₃ carbide core [11]. They also reported that the transition from M₇C₃ to M₂₃C₆ increased with raising destabilization temperature [11]. The direct transformation from M₇C₃ to M₂₃C₆ was not readily observed in the present study. Although the transformation of carbides was not directly observed, the formation of M₂₃C₆ is expected to be caused by destabilization of austenite, and the reaction between matrix and M₇C₃ is expressed with Equation (2) as follows:

M₇C₃ + γ → M₂₃C₆ + γ*

(2)

where γ* is austenite with lower alloy content than that of the original matrix γ [4,8].

The interface of M₇C₃ had a relatively smooth interface and a narrow precipitation-free zone (PFZ) with fine M₂₃C₆ particles at the outer region of PFZ in the conventionally heat-treated alloys (in Figure 6, Figure 7, Figure 8 and Figure 9). This means that the reaction of Equation (2) initiated at the interface of the carbide was directly contacted with a matrix-like peritectoid reaction [11]. Thus, the M₂₃C₆ carbides appeared along the PFZ in the vicinity of eutectic M₇C₃ reaction in Equation (2) due to dissolution of M₇C₃ into the enveloped austenite PFZ at destabilizing temperature, followed by precipitation from the saturated austenite (PFZ) (Figure 11). The formation of a small gap of PFZ, which meant M₇C₃ at the interface dissolved into the matrix (Equation (2) reaction), led to the precipitation of M₂₃C₆ from the saturated PFZ.

It is known that the eutectic M₇C₃/austenite interface does not act as a preferential heterogeneous site for M₂₃C₆ precipitation due to the following two reasons [9,10,11]: (1) formation of M₂₃C₆ at the shell of M₇C₃ along eutectic structure is perhaps a peritectoid-type reaction consuming both M₇C₃ and austenite and (2) the formation of M₂₃C₆ may not result only from the transformation of M₇C₃ to M₂₃C₆ but also from secondary precipitation from the adjacent matrix. As shown in Figure 6, Figure 7, Figure 8 and Figure 9, the formation or precipitation of M₂₃C₆ in the dendritic regions and along the periphery of M₇C₃ occurred in all conventionally heat-treated alloys. However, precipitation of M₂₃C₆ in the eutectic austenite scarcely occurred in the conventionally heat-treated alloys in the present study.

The DA-treated alloys (2124, 2127, and 2427) with a relatively high fraction of the primarily solidified dendrites had both the reaction product M₂₃C₆ (between M₇C₃ and austenite) and destabilized product (from saturated dendritic and eutectic austenites). The M₂₃C₆ precipitation in the dendritic austenite occurred due to the destabilization of austenite, which led the solute elements to emerge to release saturation. However, some portion of M₂₃C₆ in the eutectic austenite is supposed to occur with the reaction between the matrix and M₇C₃ (Equation (2)) and destabilization (less than that in the dendritic austenite). As shown in Figure 1, the fraction M₇C₃ suddenly decreased at a high temperature below the freezing temperature and that of M₂₃C₆ took over the portion of M₇C₃ reduction. Though the inverse proportion between the fractions of stable M₂₃C₆ and metastable M₇C₃ might explain the partial transformation or replacement, as mentioned in the previous reports [11,12], the M₂₃C₆ precipitation began to occur in the vicinity or at the interface of M₇C₃, rather than in its core, at the early stage of DA (less than 60 min at each DA treatment), as shown in Figure 10, Figure 11, Figure 12, Figure 13, Figure 14, Figure 15, Figure 16 and Figure 17. This phenomenon indicates that some portion of the precipitation was caused by Equation (2) reaction between M₇C₃ and austenite. In particular, precipitates around M₇C₃ in 240 min or 300 min DA exposure of each alloy (in Figure 12, Figure 13, Figure 14, Figure 15, Figure 16 and Figure 17) were very fine, compared with coarse particles in the dendritic austenite that grew after destabilization. In fact, the total amount of M₂₃C₆ was composed of the precipitation amount in both dendritic and eutectic austenites resulting from destabilization, and the reaction product of Equation (2). M₂₃C₆ precipitation preferentially occurred in the dendritic austenite in the conventional heat treatment; however, it also occurred in the eutectic austenite during DA treatment, which might be due to the solubility limit in austenite depending on temperature. In gray iron [2], graphitization occurs due to the solubility limit of C in austenite depending on temperature and C content. Thus, the experimental alloys may have their own solubility limit depending on the temperature and nature of austenite. As mentioned earlier, the eutectic austenite may have relatively low contents of Cr and C, in contrast to those in the dendritic austenite, due to the formation of M₇C₃ eutectic carbide during the eutectic reaction (which is expected to consume Cr and C). Thus, destabilization of the eutectic austenite was less active during the conventional heat treatment at 1065 °C; however, destabilization of eutectic austenite might be active even at relatively low temperatures in DA treatments.

The PFZ along the periphery of M₇C₃ in the conventionally heat-treated alloys lacks the alloying elements C and Cr, as reported earlier [3,4]. The dissolution of M₇C₃ would occur preferentially at the interface during DA treatment, at the temperature at which maximum M₂₃C₆ precipitation occurred and M₇C₃ was unstable or metastable, which might then supply alloying elements to the PFZ. As time elapsed at high temperature (during DA treatment), the concentration of the alloying elements in the PFZ was saturated, followed by precipitation of M₂₃C₆. Therefore, M₂₃C₆ precipitation along the periphery of M₇C₃ was very active during DA treatment due to the stability of M₇C₃ at the DA treatment temperature.

Figure 18a,b display the interface morphology of M₇C₃ in the DA-treated specimens. The reaction between M₇C₃ and austenite occurred regardless of the alloy composition. As already shown in Figure 11, the interface had a very small PFZ gap between M₇C₃ and M₂₃C₆ fine precipitates. The small PFZ gap presumably formed by the dissolution of M₇C₃ at the very interface, followed by M₂₃C₆ precipitation from the saturated PFZ gap. Then, the M₂₃C₆ precipitation depleted the solute elements in the PFZ, and thus, the reaction between M₇C₃ and PFZ (austenite) progressively continued.

In order to observe the progressive reaction and PFZ movement with exposure time, vacuum heat exposure on the DA-treated specimens was conducted at each DA treatment temperature for 60 more minutes. The DA-treated (for 240 min) 2124 and 2127 alloys were subjected to 60 more minutes under 1 × 10⁻⁵ mmHg vacuum, followed by Ar gas fan quenching. The microstructural observations were carried out at the same area of each specimen before and after the vacuum heat exposure.

As displayed in Figure 18, the size of M₇C₃ became smaller, and the interface became wavy. PFZ gap did not show a difference between before and after vacuum heat exposure. Increased precipitation and growth of M₂₃C₆ occurred with heat exposure. If a direct transition from M₇C₃ to M₂₃C₆ occurred at the interface of M₇C₃, the remaining M7C3 particles could not leave PFZ with heat exposure. After heat exposure, the PFZ gap was similar to that of before heat exposure, though the size of the M₇C₃ particles decreased, and the interface became wavy. From the results of the heat exposure under vacuum, it was possible to confirm that the eutectoid reaction in Equation (2) occurred between eutectic M₇C₃ carbide and austenite.

Therefore, M₂₃C₆ precipitation occurred in all areas of DA-treated specimens, unlike little or no precipitation that occurred in the eutectic region of the conventionally heat-treated specimens. In other words, most of the M₂₃C₆ precipitation was attributed to the destabilization of the matrix during conventional heat treatment; thus, the precipitation in the eutectic austenite was not active, as displayed in Figure 5. However, M₂₃C₆ precipitation during DA treatment occurred in all regions (including both eutectic and dendritic austenites) of the alloys because of Equation (2) reaction and solubility difference in austenite depending on temperature.

4. Summary

M₂₃C₆ precipitation behavior in high-Cr white iron with a composition range of 2.1~2.9 wt.% C and 24.0~27.0 wt.% Cr was studied and reviewed with a variety of chemical compositions. The summary is as follows:

• It was found that destabilization of austenite during conventional heat treatment releases saturated solute elements, C and Cr, to form M₂₃C₆ in the dendrite austenite; however, little M₂₃C₆ precipitation occurred in eutectic austenite.
• The amount of M₂₃C₆ precipitation during destabilization was closely related to that of the primarily solidified dendrites, which means that it depended on chemical composition.
• The alloy whose solidification began with delta ferrite had a relatively small fraction of M₇C₃, but precipitation of M₂₃C₆ within dendrite was very active.
• M₂₃C₆ precipitation was caused by the following two phenomena:

  (1)

  Destabilization of dendritic austenite, which released saturated solute elements, C and Cr, to form M₂₃C₆ in the austenitic matrix. Additionally, destabilization of eutectic austenite occurred at a relatively low temperature.

  (2)

  The reaction between eutectic M₇C₃ carbide and austenite with the following reaction:

  M₇C₃ + γ → M₂₃C₆ + γ*

• Direct aging at the temperature of maximum M₂₃C₆ fraction (at which M₇C₃ is expected to be metastable) developed uniform precipitation of M₂₃C₆ across the materials. M₂₃C₆ precipitation during conventional heat treatment preferentially occurred within the primarily solidified dendrites and along the periphery of M₇C₃, whereas direct aging resulted in both dendrite and eutectic austenites.
• During direct aging, both the destabilization in dendritic and eutectic austenites and the reaction between M₇C₃ and austenite simultaneously occurred.

5. Future Work

The high-Cr white iron was applied to the wear-resistant components. The wear properties of high-Cr white iron were closely related to the types and the fractions of carbides and the applied stress condition. Relatively simple DA heat treatment developed a high fraction of M₂₃C₆ carbide, compared with that of the conventional heat treatment. The existence of a high fraction of relatively fine M₂₃C₆ particles is expected to improve the wear property under relatively low-stress conditions; however, that of M₇C₃ would be effective under high-stress conditions. The authors plan to carry out a study on the effect of DA treatment temperature on microstructural evolution and wear property under various stress conditions in the future. The results may contribute to the development of proper microstructure for a specified stress condition.