Carbide transformation behaviors of a Cr–Mo–V secondary hardening steel during over-ageing

The transformation of carbides in a 1.9Cr-1.4Mo-0.3 V secondary hardening steel that was subjected to over-ageing at 600 °C–700 °C has been investigated. The carbides were characterized using scanning electron microscope (SEM), x-ray diffraction (XRD), inductively coupled plasma-atomic emission spectrometry (ICP-AES), and transmission electron microscopy (TEM) preformed on carbon replicas. The results indicate that MC, M2C, and M3C were formed during over-ageing from 600 to 700 °C, whereas M7C3, and M23C6 started to be formed at 650 and 700 °C, respectively. In addition, the co-existence of hexagonal and orthorhombic M7C3 structures in a carbide particle was firstly observed. M3C was transformed to other carbides, and the formation of both M2C and M23C6 may follow the ‘separate nucleation’ mechanism, whereas M7C3 was transformed from M3C via the ‘in situ nucleation’ mechanism. The crystallographic orientation relationships between the in situ transformed M7C3 and M3C are ( 11 2 ¯ ) M 3 C / / ( 3 3 ¯ 0 1 ¯ ) M 7 C 3 and [ 312 ] M 3 C / / [ 10 1 ¯ 3 ] M 7 C 3 .


Introduction
Cr-Mo-V steels are widely used in the hot mold industry, due to their excellent room temperature and high temperature strength [1][2][3], which is strongly correlated with the precipitation of carbides in martensite [4][5][6][7].
The behavior of alloy steels following carbide transformation during tempering has been widely studied, and has been reported for transformations such as M 3 C to MC [8], M 3 C to M 2 C [9,10], M 3 C to M 7 C 3 [11][12][13], M 3 C to M 23 C 6 [14,15], M 7 C 3 to M 23 C 6 [16,17], M 7 C 3 to M 6 C [18], MC to M 2 C [19], M 2 C to M 6 C [20] and M 23 C 6 to M 6 C [21]. Tsai and Yang [11] observed the transformation of M 3 C to M 7 C 3 in the martensitic region, resulting from the 'in situ nucleation' mechanism. Sung et al [17] concluded that M 7 C 3 carbides are directly formed from the supersaturated matrix of solutes, or through the dissolution of pre-formed M 3 C carbides, into a matrix and the subsequent nucleation reaction. For these reasons, it has been suggested that some carbides may be transformed to produce different carbides, via two basic mechanisms [22,23]: (1) 'separation nucleation', in which the pre-existing carbides dissolve firstly into the matrix, and the new carbide then nucleates and grows. In general, the new alloy carbides nucleate separately at new sites, primarily on dislocations, grain boundaries and sub-boundaries; (2) 'in situ nucleation', during which the new carbide nucleates directly from the pre-existing carbide, and grows at the expense of the mother carbide.
In order to gain an in-depth understanding of the carbide transformation mechanism, it is essential to characterize the crystallographic orientation relationships (CORs) between the carbides. However, relatively few studies have been published concerning the CORs between in situ transformed carbides. This may be due to the difficulties encountered in (1) capturing the in situ transformed complex carbides among the tens of thousands of carbides in the specimen, and (2) precisely indexing the complex selected area diffraction patterns produced by the in situ transformed carbides. The CORs between some in situ transformed carbides has been characterized and documented, such as ( ) 4151 M C 7 [21]. However, the CORs between M 3 C and its subsequent in situ transformed carbides are rarely reported, especially in Cr-Ni-Mo-V secondary hardening steels during an over-ageing stage. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI.
The aim of this study is to reveal the carbide transformation mechanism occurring in a 1.9Cr-1.4Mo-0.3 V secondary hardening steel during an over-ageing stage, and in particular, to characterize the CORs between the in situ transformation of carbides from M 3 C to M 7 C 3 .

Experimental procedure
In the present study, a 1.9Cr-1.4Mo-0.3 V secondary hardening steel was used. The composition of the steel is listed in table 1. The blank samples (Φ36 mm×L70 mm) were cut from an annealed steel bar along the rolling direction. All samples were initially heat-treated at 950°C for 1 h (austenitizing ), followed by an oil cooling process, and were then tempered at 600, 625, 650, 675 and 700°C for 2 h, respectively (the over-ageing stage).
To know the distribution of carbides, the sample was cut into 10×10×10 mm, and its cross section was polished and then was etched with 4% nitric acid alcohol. And then, the microstructure was observed by JEOM-7800 field emission scanning electron microscope.
For qualitative and quantitative analysis of carbides, the over-ageing samples were machined to produce cylindrical specimens with a diameter of 8 mm and a length of 70 mm. They were subsequently dissolved in a 1% tetramethylammonium chloride +10% acetylacetone methanol solution, to achieve precipitation using an electrolytic extraction technique. After washing in a 10 g l −1 citric acid ethanol solution, the extracted precipitation was filtered using a microporous filter film, thoroughly dried at 70°C, and weighed with an electronic balance accurate to ±0.1 mg. To identify the precipitation, the extracted powders was analyzed with a Panacow X'pert MPD x-ray diffractometer using Cu Kα radiation. In order to separate the different carbides present in the precipitation, the latter was immersed in a 6% (V/V) H 2 SO 4 +20% (V/V) H 2 O 2 aqueous solution in a bath heated with boiling water, for 1∼1.5 h. During this process, hydrogen peroxide was added to dissolve the M 3 C, MC and M 2 C; and then to separate the M 7 C 3 . Another step, the precipitation was placed in a 5∼10% (V/V) HCl ethanol solution for 0.5∼1 h, allowing the M 3 C phase to be dissolved, and the MC and M 2 C to be separated. The separation methods were referred to the literature [24][25][26]. In order to quantify the precipitation, it was washed three times, firstly with a 10 g l −1 citric acid ethanol solution, and then with a 10 g l −1 citric acid solution. Finally, it was washed with distilled water. The filter film and precipitation were transferred to a Teflon beaker with 10 ml HCl, 1 ml HNO 3 and 1 ml HF, which were added to dissolve the sample, and then diluted with water in a 100 ml Teflon volumetric flask at constant volume. The content of each element was determined using an inductively coupled plasma-atomic emission spectrometer (ICP-AES).
To further characterize the precipitation (morphology, size, etc), the precipitation extracted by carbon replicas was used to prepare the TEM sample. Firstly, a thin carbon film was evaporated onto the surface of the etched steel metallographic sample, using SBC-2 multi-functional sample surface treatment instrument. The carbon film was then stripped from the surface and the precipitation was extracted from the surface region of the martensite matrix by means of deep etching with a 10% nitrate acid alcohol. Finally, it was placed on a 300-mesh Cu grid, and was examined in a JEM-2100F TEM operating at 200 kV. JMatPro 7.0 was used to calculate the kinetics of the carbide transformation, which is based on the classical Johnson-Mehl-Avrami theory with an adaptation that allows the morphology of the precipitate to be considered as well as specifics associated with potential nucleant sites [27]. The carbide evolution holding at 600, 650 and 700°C, respectively, after austenitizing at 950°C, was computed by JMatPro 7.0.

Results and discussion
3.1. Carbide identification and morphologies Figure 1 shows the SEM images of the samples at different over-aging temperatures. The microstructures consist of tempered martensite matrix and carbides in globular, needle-like and rod-like shapes. The globular and rodlike carbides are mainly distributed at the martensite lath and the grain boundaries, while the needle-like carbides are distributed between the lath and the grain boundaries. Figure 1(f) demonstrates the EDS energy spectra of the red-marked carbide in figure 1(a). This result indicates that the globular carbide is a V-rich phase, which is likely to be MC carbide. Figure 2 shows the XRD patterns of the powders extracted from the samples during the over-ageing stage, between 600 and 700°C. Orthorhombic M 3 C, hexagonal M 2 C, and face-centered cubic MC were detected throughout the full duration of the over-ageing stage tests, whereas hexagonal M 7 C 3 was detected only at temperatures higher than 625°C. Figure 3 illustrates typical TEM images and SAED patterns of carbides at 600, 675 and 700°C, respectively. The results indicate that the MC has a globular shape, the M 2 C is rod-like, the M 3 C is needle-like, and the M 7 C 3 is polygonal with well-defined edges. Figure 4 shows the TEM and SAED pattern for the rod-like M 23 C 6 at 700°C.     The M 23 C 6 carbide was not detected by XRD analysis, which may have been due to its low content. It is interesting to note that orthorhombic M 7 C 3 was also observed by TEM and SAED, and this is further discussed in the following section. The crystallographic characteristics of these carbides, derived from XRD, TEM and ICP-AES, are summarized in table 2.   Figure 5 plots the ICP-AES characterized content variations of three different types of carbide, as a function of tempering temperature. The weight percentage f of the carbides extracted from the specimens was calculated using the following formula:

Carbide precipitation sequence
where m c is the weight of each specific carbide, and m o and Mr are the weights of the specimen before and after electrolytic extraction, respectively. As it is difficult to distinguish MC from M 2 C using the electrolytic extraction technique, the sum of the weight percentages of MC plus M 2 C is shown (red curve). These results show that at temperatures ranging between 600 and 700°C, M 3 C decreases from 2.42 wt% to 0.56 wt%, whereas the sum of MC plus M 2 C increases from 0.76 wt% to 1.18 wt%. M 7 C 3 is initially detected at 650°C, and its weight content then increases with increasing temperature. Figures 6(a)-(c) plots the weight content variations of various metals detected in the M 3 C, M 2 C and MC, and M 7 C 3 carbides, respectively, following ICP-AES characterization. These results show that as the temperature increases, the weight content of all detected metals (i.e. Fe, Cr, Mo, V, Mn) and C decreases in M 3 C, implying that this carbide is progressively dissolved. Under the same conditions, the weight content of Fe, Cr and Mo in M 2 C and MC increases, whereas the weight content of V in these carbides remains constant. This implies that the precipitation of M 2 C increases, whereas MC remains insensitive to tempering temperature. M 7 C 3 is rich in Cr and Fe, and contains only small relative amounts of Mo, V and Mn. The weight content of all elements in M 7 C 3 also increases with increasing temperature. Figure 6(d) plots the combined weight content variation of these carbides as a function of tempering temperature. It can be seen that the total carbide content varies only slightly over the range 600°C-675°C, and declines somewhat when the temperature increases to 700°C. This reveals that at 700°C, a small amount of these carbides has been dissolved into the martensite matrix, leading to an increased alloy content in the matrix.
The results obtained using TEM and ICP-AES thus lead to the conclusion that, as the tempering temperature increases during the over-ageing stage, M 3 C is gradually converted to M 2 C, M 7 C 3 and trace quantities of M 23 C 6 . The carbide evolution holding at 600, 650 and 700°C, respectively, after austenitizing at 950°C, was computed by JMatPro 7.0, as shown in figure 7. In the temperature range, the M 3 C dissolves gradually with increased precipitation of M 2 C and no precipitation of MC as the holding time progresses, which is slightly different from the experimental result. The precipitation of M 7 C 3 begins at approximately 650°C, and becomes clearly visible at 700°C, which is consistent with the experimental measurements (figures 5 and 6). A small quantity of M 23 C 6 was detected at 700°C, thus further confirming the existence of M 23 C 6 at high temperatures.
Following the results of theoretical calculations and experiments on the carbide precipitation, our proposed carbide transformation sequence during the 600°C-700°C over-ageing stage is: M 3 C+M 2 C+MC→M 3 C+ M 2 C+MC+M 7 C 3 →M 3 C+M 2 C+MC+M 7 C 3 +M 23 C 6 . Many studies [3,7] have revealed that carbide transformation in Cr-Mo steels during the over-ageing stage is predominantly correlated with the Cr/Mo ratio, with the preferential formation of M 2 C, M 7 C 3 and M 23 C 6 corresponding to low, moderate and high Cr/Mo ratios, respectively.

Carbide transformation mechanism
It is well-known that carbide transformation takes place either through separate nucleation or by way of in situ nucleation. The indexing SAED patterns shown in figure 8(d) reveal that the diffraction spots generated by different phases are entangled, which is completely different to the ordered pattern to be expected with a single phase. This shows that the transformation of M 3 C and M 7 C 3 took place according to an in situ nucleation mechanism. The observation of this in situ nucleation was also found in other carbide particles, as shown in figure 9. The crystallographic orientation relationships between the M 3 C and M 7 C 3 obtained from the SAED diffraction pattern were identified as follows: In addition, the co-existence of hexagonal and orthorhombic M 7 C 3 structures was observed in a carbide particle. Although hexagonal and orthorhombic M 7 C 3 structures have been reported in the literature [28], this is the first time that both structures have been detected in a single carbide particle transformed by in situ nucleation. The nucleation of carbides at various types of boundary is to be expected since these are energetically