Effect of cerium addition on microstructure and mechanical properties of high-strength Fe85Cr4Mo8V2C1 cast steel

Materials Science & Engineering A 674 (2016) 366–374

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Effect of cerium addition on microstructure and mechanical properties of high-strength Fe85Cr4Mo8V2C1 cast steel J. Hufenbach a,n, A. Helth a, M.-H. Lee b, H. Wendrock a, L. Giebeler a, C.-Y. Choe b, K.-H. Kim b, U. Kühn a, T.-S. Kim b, J. Eckert a,c,1 a

IFW Dresden, Institute for Complex Materials, P.O. Box 270116, D-01171 Dresden, Germany Korea Institute of Industrial Technology, Gaetbeol-ro 156, Yeonsu-gu, Incheon 406-840, Korea c TU Dresden, Institute of Materials Science, D-01062 Dresden, Germany b

art ic l e i nf o

a b s t r a c t

Article history: Received 27 June 2016 Received in revised form 26 July 2016 Accepted 27 July 2016 Available online 3 August 2016

This work presents an investigation on the influence of rare earth additions (Ce) on the microstructure and mechanical properties of a cast Fe85Cr4Mo8V2C1 (element contents in wt%) tool steel. The applied relatively high solidification rate during the casting process promotes the formation of non-equilibrium phases such as martensite, retained austenite as well as a fine network-like structure of complex carbides. This combination of phases and their morphology results in excellent mechanical properties already in the as-cast state. Cerium additions induce a change in phase formation and resulting mechanical properties. Besides morphological and quantitative changes of the main constituent phases, novel carbo-oxide and carbide phases are formed. To investigate this microstructural phenomenon, X-ray diffraction (XRD), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) combined with energy dispersive X-ray spectroscopy (EDX) were applied. Altogether, the addition of small amounts of the rare earth element cerium together with a tailored casting process results in enhanced mechanical properties compared to the Fe85Cr4Mo8V2C1 alloy and offers new possibilities to obtain high-strength and simultaneously adequate ductile cast steels for advanced tool design. & 2016 Elsevier B.V. All rights reserved.

Keywords: Steel Casting Ce addition Microstructure Mechanical characterisation

1. Introduction The increasing requirements for processing novel high-performance materials as well as general demands for longer service life and more efficient processes are the greatest challenges for tool making at present and in the future. Thereby the application of further improved high-performance tool steels plays a key role. Cast high carbon Fe-(Cr)-(W)-Mo-V alloys show a large potential to replace wrought alloys for various applications in tool manufacturing due to their high hardness, wear-resistance and compressive strength [1–6]. Furthermore, the production of nearnet-shape tools made of cast steels is an increasing economical alternative to forged or rolled steels, because no long-term, multistage manufacturing process is required. The casting technology additionally allows a high flexibility concerning the used materials n

Corresponding author. E-mail address: [email protected] (J. Hufenbach). 1 Present address: Erich Schmid Institute of Materials Science, Austrian Academy of Sciences and Department Materials Physics, Montanuniversität Leoben, Jahnstraße 12, A-8700 Leoben, Austria. http://dx.doi.org/10.1016/j.msea.2016.07.114 0921-5093/& 2016 Elsevier B.V. All rights reserved.

as well as the product design. By applying near-net-shape casting, subsequent mechanical machining can be reduced or completely avoided. Furthermore, casting enables an adjustment of the desired microstructure and the resulting properties by tailored temperature control [4,7,8]. By applying e.g. high cooling rates a fine microstructure and metastable phases can be adjusted, which can also be very beneficial for further improving the mechanical properties of tool steels [3,4,6]. Nevertheless, steels produced by casting processes may also show pores and coarse eutectic carbide structures [5], which can promote crack growth and early failure, and should therefore be reduced by appropriate methods. The effect of rare earth (RE) metals in steels as deoxidizer and desulphuriser is well known from literature [9,10] as well as the related influence on the solidification process [1,11]. It has been reported that RE elements like cerium, lanthanum and yttrium can form highly stable oxides, oxy-sulphides and sulphides [10,12–14]. Due to the very low Gibb´s energies of formation of these compounds even at high temperatures, the RE elements can immediately react with O and S when added to the steel melt [15]. These compounds precipitate as solid particles in the melt due to their high melting temperatures and trigger grain refinement [10].

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There are only a few reports on the improvement of mechanical properties and related changes in microstructure due to RE addition [1,11,12,16–18]. It was shown that the grain size of tool steel could be refined, which may result in increased strength, and that the morphology of eutectic carbides could be changed from its network-like structure to granular carbides (especially after subsequent heat-treatment) resulting in enhanced impact and fracture toughness [10,15,18]. Furthermore, the addition of Ce minimises the segregation of W and Mo in tool steels and may decrease the amount of eutectic carbides as well as refine carbide lamellae [1]. The influence of RE metals on the mechanical properties of tool steels is, however, still controversially discussed. In general, the reason for the enhancement of the mechanical properties points to morphological changes of the phase constituents and formation of various precipitations [1,11,12,16,17]. The present Fe85Cr4Mo8V2C1 (wt%) cast alloy was produced at relatively high solidification rates (about 10–70 K/s [3]) and pure preparation conditions. Thereby a microstructure composed of martensite, retained austenite and a carbide network is already obtained in the as-cast state [3]. This combination of phases as well as their morphology results in excellent mechanical properties, especially a high strength combined with an adequate fracture strain under compressive loading [3]. The strength is thereby already comparable with other cast tool steels after a long-term, cost-intensive heat-treatment process [5], so that the combination of a special casting process with a tailored alloy is very promising. The aim of the present work is to study the effect of small Ce additions on the microstructure and mechanical properties of the Fe85Cr4Mo8V2C1 alloy. For further improvement of the mechanical properties and simultaneous reduction of pores in the cast ingot, small amounts of cerium with concentrations of 0.03, 0.1, 0.3 wt%, respectively, were added to the alloy. Intensive investigations regarding the morphological changes of the microstructural constituents as well as the influence of the changing phase fractions on the hardness and mechanical properties under compressive and tensile loading were performed to allow a profound understanding on the impact of cerium additions on microstructure development and deformation characteristics.

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spectroscopy (ICP-OES; IRIS Intrepid II XUV, Thermo Fisher Scientific) for the metal concentrations. For investigating the microstructure of the as-cast samples, optical microscopy (OM; Epiphot 300, Nikon) as well as scanning electron microscopy (SEM; Leo 1530 Gemini, Zeiss) in combination with energy dispersive X-ray spectroscopy (EDX; X-Flash, Bruker) and wavelength dispersive X-ray spectroscopy (WDX; Microspec2, Oxford Instruments) was applied. High resolution analyses of the oxidic and carbidic compounds were carried out by transmission electron microscopy (TEM; Tecnai G2 F20, FEI). For the OM observations, the samples were etched with Beraha I reagent (stock solution: 1 mL H2O, 200 mL HCL, 24 g (NH4)HF2; 100 mL stock solution, 1 g K2S2O5) after grinding and polishing. The SEM overview images were performed on Nital etched samples (10 mL HNO3, 90 mL C2H5OH), whereby the carbide morphology was examined on deep-etched samples (5 g FeCl3, 10 mL HNO3, 3 mL HCl, 87 mL C2H5OH). The fracture surfaces of the compression and tensile test samples were also investigated by SEM. For the TEM investigations, a cross-sectional TEM specimen out of the Ce0.3 alloy was prepared by focused ion beam (FIB). The prepared TEM lamella was then attached to a Cu-grid in FIB. X-ray diffraction (XRD; STOE Stadi P, Mo Kα1 radiation, Mythen 1K detector, Ge(111) monochromator) combined with Rietveld analysis [21] by using Fullprof [22] was applied for the qualitative and quantitative phase analysis. The quasi-static compression tests (Shimadzu AG-250KNI) were conducted at a strain rate of 10
3 s
1 by using grinded cylindrical and coplanar samples with dimensions of 3 mm diameter and 6 mm length. The quasi-static tensile tests were performed at a strain rate of 10
3 s
1 using an Instron 8562 testing device. Thereby rectangular dog-bone-shaped samples according to DIN 50125 were prepared from the cast ingots. For the determination of the material parameters from the compression and tensile tests, at least three samples were tested. Based on this, the arithmetic averages and standard deviations were calculated. In addition to the compression and tensile tests, Rockwell macro-hardness measurements (CV Instruments) were performed at a test load of 1471 N. Ten measurements were used for the ascertainment of arithmetic averages and standard deviations, respectively.

2. Experimental The Fe85
xCr4Mo8V2C1Cex (x ¼0.03, 0.1, 0.3 wt%) alloys [20] were prepared by induction melting (Balzers) of the pure elements in a ceramic crucible under argon atmosphere. At a temperature of about 1823 K, the melt was cast into a CuZr mould to obtain ingots with dimensions of (70  120  14) mm3. To achieve the desired microstructure composed of martensite, retained austenite and carbides, an average solidification rate of about 10–70 K/s has to be realised [3]. The nominal chemical compositions of the four cast ingots (see Table 1) were confirmed by carrier gas hot extraction (CGHE; EMIA 820V, Horiba) for determining the carbon concentrations and by inductively coupled plasma optical emission

3. Results and discussion 3.1. Microstructure Fig. 1a presents an optical micrograph of the dendritic microstructure of the as-cast Fe85Cr4Mo8V2C1 base alloy after etching with Beraha I solution. Previous studies on this alloy have revealed that the dendrites are mainly composed of martensite appearing as blue needles in Fig. 1a. The interdendritic area contains a network-like structure of bright V-rich and Mo-rich carbides surrounded by light yellow coloured austenite (Fig. 1a) [3]. This

Table 1 Nominal and experimental chemical composition (in wt%) of the investigated alloys. Alloy

Chemical composition

Fe (wt%)

Cr (wt%)

Mo (wt%)

V (wt%)

C (wt%)

Ce (wt%)

FeCrMoVC

Nominal Experimental Nominal Experimental Nominal Experimental Nominal Experimental

85 84.77 84.97 85.60 84.9 85.02 84.7 84.48

4 3.93 4 4.03 4 4 4 3.98

8 7.98 8 7.93 8 7.95 8 7.94

2 2.01 2 2.01 2 1.99 2 1.98

1 1.01 1 1.0 1 0.98 1 1.0

– – 0.03 0.029 0.1 0.11 0.3 0.27

Ce0.03 Ce0.1 Ce0.3

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Fig. 1. (a) OM image of the dendritic microstructure of the cast Fe85Cr4Mo8V2C1 alloy etched with Beraha I, (b–e) SEM images of the (b) Fe85Cr4Mo8V2C1 base alloy and its modifications (c) Ce0.03, (d) Ce0.1 and (e) Ce0.3; white circles mark cavities. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).

Fig. 2. (a) SEM image of the carbide network of the cast Fe85Cr4Mo8V2C1 alloy and (b) the modification Ce0.3; both after deep-etching.

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Fig. 3. (a) SEM image of the Ce0.3 alloy displaying the carbide network within the matrix. Related WDX mapping displaying the distribution of (b) O as well as EDX mappings showing the distribution of (c) Ce, (d) C, (e) Cr, (f) Fe, (g) Mo and (h) V.

complex microstructure composed of martensite, austenite and carbides is also reflected in the OM images of the cerium-containing alloy modifications. Fig. 1b–e display SEM micrographs of the as-cast Fe85Cr4Mo8V2C1 base alloy and alloy modifications with 0.03, 0.1 and 0.3 wt% cerium after etching with Nital solution. The SEM overview images of the alloy modifications (Fig. 1c–e) indicate that the addition of cerium affects the carbide morphology. In comparison with the FeCrMoVC base alloy, changes from lamellar, complex branched carbides to more compact, rounded carbides are observed. In addition, necking and interruption occur at the junctions of the carbide network with increasing Ce content. This is confirmed by the SEM micrographs of deep-etched samples shown in Fig. 2a and b. In both samples characteristic compact as well as branched lamellar carbides appear, which were determined as V-rich MC-type with M ¼V, Mo, (Cr) and Mo-rich M2C-type (M ¼Mo, V, Cr) carbides in a previous study, respectively [3] (see Fig. 3). After deep-etching it can be seen that the carbides become finer and more compact with increasing Ce additions when compared to the carbides in the FeCrMoVC base alloy (Fig. 2a and b). This phenomenon accords to the findings described by Wang et al. [11], Fu et al. [1] and Yang et al. [23]. Furthermore, several round-shaped precipitates are clearly visible in the SEM micrographs (shining white coloured particles in Fig. 3a) of the 0.3 wt% Ce-containing alloy modification. For analysing the element distribution in this phase EDX and WDX examinations were performed (Fig. 3b–h). Thereby Ce, O and C could be detected in all examined precipitates. Additionally, small amounts of Cr, V and Mo were verified in this phase. These findings indicate that the round-shaped precipitates in Fig. 3a can be Ce-rich carbo-oxides. The formation of such complex oxides caused by rare earth additions in steels with similar morphologies is also described in literature [10,12,19]. It is known that RE elements have a high affinity to oxygen and can directly react in the melt forming high-melting oxides due to the very low Gibb's energies of formation of these compounds [10,15]. These oxides precipitate as solid particles in the melt and can act as heterogeneous nuclei of primary austenite under certain conditions [10,16,17]. With increasing nucleation of primary austenite, it is assumed that the remaining melt is separated off in pools before the eutectic reaction begins. Therefore, the eutectic austenite mainly nucleates in the narrow areas along the primary austenite, so that fine and more isolated eutectic carbides are formed [16,17]. Those carbides were also observed in the examined Ce

modifications as described above. Besides sample Ce0.3, the same particles were also found in the samples Ce0.1 and Ce0.03 (not shown), and the volume fraction of this assumed Ce-rich carbo-oxides seems to increase with increasing Ce content, similar as described by Fu et al. [1] and Gao et al. [12]. Furthermore, EDX analyses confirm a slight enrichment of Ce in the V-rich MC-type and the Mo-rich M2C-type mixed carbides. The details of the Ce-rich carbo-oxides are investigated by using TEM. Fig. 4a–f show a TEM bright-field image displaying the Ce-rich carbo-oxides (Fig. 4a) with the corresponding EDX mapping results for Fe, Ce, O, Mo and C (Fig. 4b–f). The elemental maps of V and Cr are not shown here due to the superposition of the V Kα signal with the strong Ce Lα signal and the Cr Kα signal with Ce Lβ. As it can be seen in Fig. 4a, two different areas can be observed in the recorded TEM image (marked as I and II). The chemical analysis results indicate higher concentrations of Ce, O, and C in area II (Fig. 4c, d and f) than in area I and the Fe-rich matrix. Based on the chemical investigation, the structural investigation was carried out with the help of electron diffraction technique in TEM. The insets in Fig. 4a respectively show the selected area electron diffraction (SAED) pattern recorded from area I and II. The recorded SAED patterns were indexed based on possible Ce-O-C phases expected from TEM-EDX analysis [24–30]. The best fit solution and related crystallographic information used for indexing the electron diffraction patterns are listed in Table 2. Thereby the indexing results show that area I consists of CeO2 (cubic, Fm-3m) nanocrystallites, whereas area II is indexed as Ce2O3 (trigonal P3m1). It can therefore be concluded that the Ce0.3 alloy contains Ce-rich carbo-oxides with different crystal structures of CeO2 and Ce2O3. Volume fractions and crystal structure types of the phases were examined by XRD. The corresponding patterns are shown in Fig. 5a and the results are summarised in Table 3. The measurements confirm the conclusions from the microscopic examinations. All investigated alloy modifications are composed of martensite, austenite and carbides. The two main phases austenite and martensite are indexed based on structure models described by Kohlhaas et al. (martensite, SG Im-3m) [31] and Ridley and Stuart (austenite, SG Fm-3m) [32]. Furthermore, carbides of the MC- and the M2C-type are detected in all alloys. Due to EDX investigations, it is known that these carbides are complex mixed carbides (see Fig. 3b–h),

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Fig. 4. (a) TEM bright-field image of a lamella cut out through Ce-rich carbo-oxides in Ce0.3 as well as corresponding electron diffraction patterns of the cubic CeO2 precipitate in area I and of the trigonal Ce2O3 phase in area II as insets; related TEM-EDX maps of (b) Fe, (c) Ce, (d) O, (e) Mo and (f) C.

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Table 2 Crystallographic information used to index the electron diffraction patterns and best fit solutions. Phase

Crystal structure

Space group Atomic coordinates

CeO2

Cubic

Fm-3m

Ce2O3

Trigonal

P-3m1

Reference

Ce: (0, 0, 0) [25] O: (1/4, 1/4, 1/4) Ce1: (1/3, 2/3, 0.24543) [26] O1: (0, 0, 0) O2: (1/3, 2/3, 0.6471)

whereby M represents Mo, V, Cr, and Ce. However, V is the major element of the MC-type carbides, and Mo the major element in the M2C-type carbides. The structure of the MC phase is refined by a cubic structure model, SG Fm-3m, published by Pflüger et al. [33]. The M2C phase is indexed based on an orthorhombic structure model, SG Pbcn, as described by Page et al. [34]. In sample Ce0.3 reflections of an additional phase are visible. Fig. 5b shows the XRD pattern of the Ce0.3 alloy modification and of the FeCrMoVC base alloy at low 2θ angles. The pattern of the Ce0.3 sample reveals two additional reflections in comparison to the base alloy (Fig. 5b). Both are ascribed to a tetragonal phase (CeC2-type) with the space group I4/mmm[35]. Lattice parameters of the structure are estimated and presented in Table 3. The carbidic Ce-containing phase is denoted as structure type since a mixed carbide (containing Mo, Cr or V) or oxy-carbide, where some carbon may be exchanged by oxygen, is the most likely case. However, due to its low amount, this phase is not reliably quantifiable by XRD and could not be detected by other characterisation methods. The CeC2 phase may be finer and more homogeneously spread over the whole sample compared to the carbo-oxides, which can explain their detection by XRD, but not by SEM-EBSD or TEM-SAED even for the use of FIB cut-techniques. In contrast, a hit of a carbo-oxide precipitate (volume fraction below 0.1 vol%, estimated from some SEM images of (240  180) mm2) by XRD to get a more or less resolved pattern in a rotating sample would be a serendipitous observation. Table 3 compiles the structure data of the base alloy and its Cemodifications as well as the determined phase contents, revealing that the amount of carbides (mainly MC) increases upon Ce addition to the FeCrMoVC base alloy. This effect was also described by Chaus et al. for certain high-speed steels [17]. The highest amount of carbides, however, is verified for sample Ce0.03. With further increasing Ce content the amount of MC carbides slightly decreases, whereas the formation of a novel phase with CeC2 structure is promoted. Basically the same trend can be observed

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for the martensite fraction, except for alloy Ce0.1, which slightly differs. The austenite content shows an inverse trend (see Table 3). First, the austenite content decreases when adding 0.03 wt% of Ce and afterwards increases with rising Ce content, whereby the austenite fraction in Ce0.3 is similar as in the FeCrMoVC base alloy. 3.2. Mechanical properties In order to evaluate the effect of Ce on the mechanical properties of the Fe85Cr4Mo8V2C1 alloy, different investigations were carried out. The macro-hardness does not show a significant trend with respect to the Ce content of the alloys in consideration of the standard deviations. For Fe85Cr4Mo8V2C1 a hardness of 597 0.5 HRC was determined, whereby Ce0.03 shows 58.2 70.4 HRC, Ce0.1 a value of 58.6 70.7 HRC and Ce0.3 a hardness of 57.670.4 HRC. This observation leads to the conclusion, that neither the changed morphology of the carbides, nor the fluctuating phase compositions have a pronounced influence on the macro-hardness. Fig. 6a shows the engineering compression stress-strain and Fig. 6b the engineering tensile stress as function of the strain for representative Ce-modified samples. The corresponding characteristic engineering values are summarised in Table 4 and 5, respectively. Obviously, the addition of Ce and the related microstructural changes have a clear influence on the strength and deformation behaviour of the alloys under compression and tensile loading. In the compression tests, the alloy with 0.03 wt% Ce shows the highest average yield strength sy0.2 of 24087186 MPa and ultimate compression strength smax of 4421 7109 MPa (Fig. 6a). Hence, a significant improvement compared to the base alloy (sy0.2 ¼1942 738; smax ¼35367143) could be obtained. A further increase of the Ce content to 0.1 wt% results in a continuous decrease of yield and ultimate compression strength, whereby both average values of the Ce0.1 modification are still higher than the respective values of the FeCrMoVC base alloy (Table 4). Furthermore, an increasing trend of fracture strain εf and total strain εt is observed with increasing Ce content (Fig. 6a and Table 4). These trends can be explained by the morphological changes of the microstructure as well as the changing phase fractions already described above (see Section 3.1). The improvement of strength properties after the addition of 0.03 wt% Ce to the FeCrMoVC base alloy can be attributed to the lower amount of soft austenite as well as the increase of the carbide (especially MC) and martensite

Fig. 5. (a) XRD patterns of the examined alloy modifications Ce0.03, Ce0.1 and Ce0.3 in comparison with the FeCrMoVC base alloy; (b) detailed view of the XRD patterns of FeCrMoVC and Ce0.3 with indexed CeC2 phase.

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Table 3 Phase composition, space group and lattice parameters of martensite, austenite and carbides of the Fe85
xCr4Mo8V2C1Cex (x ¼0; 0.03; 0.1; 0.3) alloys determined by Rietveld analysis of XRD data. Alloy

Structure type

Space group

a [nm]

FeCrMoVC

Fe Fe0.94C0.06 VC Mo2C

Im-3m Fm-3m Fm-3m Pbcn

0.28833(4) 0.3607(1) 0.4167(6) 0.4639(7)

Fe Fe0.94C0.06 VC Mo2C

Im-3m Fm-3m Fm-3m Pbcn

0.28768(2) 0.36089(7) 0.41757(15) 0.46397(8)

Fe Fe0.94C0.06 VC Mo2C

Im-3m Fm-3m Fm-3m Pbcn

0.28773(2) 0.36135(11) 0.41745(18) 0.46423(9)

Fe Fe0.94C0.06 VC Mo2C CeC2

Im-3m Fm-3m Fm-3m Pbcn I4/mmm

0.287597(8) 0.36155(3) 0.4175(7) 0.46466(13) 0.3869(5)

Ce0.03

Ce0.1

Ce0.3

b [nm]

0.5899(11)

0.5899(2)

0.59002(2)

0.5897(3)

c [nm]

V [nm³]

Phase content [wt%]

0.5074(10)

0.02397(1) 0.0469(3) 0.073(5) 0.1389(7)

71 24 2 3

0.50843(1)

0.02307(4) 0.04700(3) 0.07381(8) 0.1391(8)

73 15 8 4

0.50825(2)

0.023838(7) 0.04718(4) 0.07275(9) 0.1392(1)

74 17 5 4

0.5084(2) 0.650(2)

0.023788(2) 0.04726(1) 0.07280(4) 0.1393(2) 0.0972(5)

69 24 3 4 –

Fig. 6. (a) Engineering compressive stress-strain curves and (b) engineering tensile stress-strain curves of the Ce-modified alloys.

Table 4 Engineering stress and strain values from room temperature compression tests of the as-cast FeCrMoVC alloy and its Ce modifications: offset yield strength sy0.2, ultimate compression strength smax, fracture strain εf, total strain εt. Alloy

ry0.2 [MPa]

rmax [MPa]

εf [%]

εt [%]

FeCrMoVC [3] Ce0.03 Ce0.1 Ce0.3

1942 7 38 24087 186 21057 85 1893 7 114

35367 143 44217 109 37627 172 3290 7 74

177 1 15.17 2.2 18.7 7 3.2 22.9 7 0.8

18.8 70.7 22.0 7 2.0 24.67 3.9 28.0 7 1.1

Table 5 Engineering stress and strain values resulting from room temperature tensile tests of the as-cast FeCrMoVC alloy and its Ce modifications: offset yield strength sy0.1, ultimate tensile strength smax, fracture strain εf, total strain εt. Alloy

ry0.1 [MPa]

rmax [MPa]

εf [%]

εt [%]

FeCrMoVC [3] Ce0.03 Ce0.1 Ce0.3

854 7 6 1036 7 20 Not determinable Not determinable

1184 769 1405 7 64 1063 7 48 10477 71

0.3 7 0.1 0.4 7 0.08 0.17 0.01 0.17 0.06

0.9 7 0.1 1.0 7 0.1 0.6 7 0.04 0.6 7 0.1

fractions (see Table 3), i.e. the increasing amount of hard phases in the alloy. Further additions of Ce lead to a significant decrease of the MC fraction from 8 wt% (Ce0.03) to 3% (Ce0.3). Even though an additional novel carbide of the CeC2 type could be detected in Ce0.1 and Ce0.3, the low total amount of carbides results in a larger amount of carbon and alloying elements in the austenite. This causes a lower martensite start and finish temperature, so that lower amounts of martensite and a higher fraction of retained austenite is formed during cooling. Besides the very small amounts of CeC2 detected with increasing Ce content, it is assumed that the rising precipitation of Ce-rich carbo-oxides with increased RE content has no dominant effect on the strength of the alloys under compressive loading. Apart from the changes in the phase fractions, one possible explanation for the observed enhanced ductility of the Ce-modified alloys may be found in the arrangement of the carbides. The complex carbide network along the austenite grain boundaries becomes increasingly interrupted upon Ce addition and crack propagation along this path is therefore hampered [36,37]. The representative tensile stress-strain curves displayed in Fig. 6b as well as the corresponding characteristic values show a similar trend concerning the yield and ultimate tensile strength with increasing Ce content. After adding 0.03 wt% of Ce to the FeCrMoVC base alloy, the

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average 0.1-yield strength sy0.1 rises from 854 MPa for the base alloy to 1036 MPa for Ce0.03. In addition, the average ultimate tensile strength smax increases from 1184 MPa for the base alloy to 1405 MPa for Ce0.03. With a further rise of the Ce content, a decreasing trend of the average ultimate strength was detected (Table 5). The 0.1% offset yield strength could not be determined for at least three tested Ce0.1 and Ce0.3 samples and therefore no values are presented in Table 5. This behaviour is mainly attributed to the changes in the phase fractions explained above. A similar decrease of the tensile strength after reaching a certain amount of rare earth addition was also described by Wang et al. [11]. The fracture strain εf and total strain εt under tensile load show nearly the same trend as for the ultimate strength. Thus, the addition of 0.03 wt% Ce leads to a comparable fracture (εf ¼0.4 70.08%) and total strain (εt ¼1.070.1) as for the base alloy (εf ¼0.3 7 0.1; εt ¼ 0.9 70.1) under consideration of the standard deviation. The addition of 0.1 and 0.3 wt% of Ce, however, results in premature fracture and, therefore, in a decrease of the ductility (Fig. 6b and Table 5). An explanation for this behaviour is found in the larger amount of cavities with increasing Ce content, as shown in Fig. 1b– e (white circles). In the FeCrMoVC alloy the M2C and MC carbides form a stable, branched network (Fig. 1a and b). Its interconnection keeps both types of carbides fixed in the surrounding austenite-martensite-matrix, so they can hardly break out upon deformation. With increasing Ce content progressive necking and interruption of the carbide network occurs and the amount of separated Ce-rich carbo-oxides increases. These (partially) isolated hard and brittle carbides as well as single oxides generally tend to break out easily during mechanical loading (see fracture surfaces in Fig. 7b and c), act as stress concentrators and are known as critical crack initiation sites in high-strength cast alloys [36]. The resulting cavities show a more detrimental effect on the deformation behaviour under tension than under compression [38]. Hence, their strong influence can dominate the failure behaviour under tensile loading. Fig. 7a–c display SEM images of the fracture surfaces of the Ce0.3 samples after compression and tensile testing. The fracture surfaces of the compression test samples of the examined alloys with Ce addition indicate shear fracture as described for the base alloy [2]. As exemplified by the fracture surface of sample Ce0.3 the evolving dimples are mainly elongated in the direction of maximum shear stress (Fig. 7a). Due to pronounced heat release caused by heavy local deformation, partially melted zones can be found at higher magnifications (Fig. 7a). This shear fracture is initiated through the formation of cavities in the vicinity of the precipitates (mainly the carbides). They grow with progressing deformation until the webs between them shear through. In contrast, the failure of all tensile test samples with Ce addition occurs as cleavage fracture perpendicular to the plane of maximum tensile stress without any clearly visible plastic deformation (Fig. 7b). This behaviour is also confirmed by the low fracture strains given in Table 5. Cracking predominantly occurs along the boundaries of the dendrites. As a consequence, the fracture surface of the Ce0.3 tensile test sample (shown in Fig. 7b) exhibits a characteristic surface relief. Broken carbides like shown in Fig. 7c are starting points for crack initiation and propagation. Similar tensile fracture features were also found for the other Cemodified alloys as well as the FeCrMoVC alloy [3]. This fracture behaviour is related to the high brittleness of the carbide network, which promotes the intergranular crack propagation [1,36].

4. Conclusions In this study the influence of Ce additions (0.03, 0.1, and 0.3 wt%) on the microstructure and the mechanical properties of a

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Fig. 7. SEM images of fractured Ce0.3 alloy specimens after (a) compression testing and (b, c) tensile testing, respectively.

cast Fe85Cr4Mo8V2C1 (wt%) tool steel was investigated. The main results are summarised as follows: All examined alloys were cast under pure preparation conditions and by realizing a high solidification rate, which promotes the formation of non-equilibrium phases such as martensite, retained austenite, and mixed carbides. For Ce contents of 0.3 wt%, two kinds of Ce-rich carbo-oxides with cubic (CeO2) and trigonal (Ce2O3) crystal structure as well as Ce carbides of a tetragonal structure type (CeC2) are found. The addition of Ce leads to a reduction of coarse carbides and initiates changes of the carbides from finely lamellar to compact, rounded-shaped. Furthermore, the carbide network shows continued necking and interruption with increasing Ce content, which was traced back to the influence of rare earth on the solidification behaviour by triggering the crystallisation of the primary austenite. Ce additions and resulting microstructural changes have a clear influence on the strength and deformation behaviour of the alloys

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under compression and tensile loading. The addition of 0.03 wt% Ce results in a significant improvement of the yield and ultimate fracture strength under compression and tensile loading compared to the FeCrMoVC base alloy. This can be attributed to the increased amount of hard phases (martensite and carbides) in the Ce modification as well as the lower amount of soft austenite. Adding higher Ce amounts (0.1 or 0.3 wt%) improves the average fracture and total strain under compressive load in comparison to FeCrMoVC. Apart from the changes in the phase fractions, this effect is mainly traced back to the changes in carbide arrangement. The complex carbide network along the austenite grain boundaries observed in the FeCrMoVC alloy becomes increasingly interrupted upon Ce addition and crack propagation along this path is therefore hampered. By adding small Ce contents to the Fe85Cr4Mo8V2C1 base alloy combined with a tailored casting process a high-strength and simultaneously adequate ductile cast steels can be obtained, which is promising for advanced tool design.

Acknowledgements The authors are grateful to H. Bußkamp, M. Frey, A. Gebert, C. Geringswald, B.-S. Kim, S.-Y. Kim, R. Keller, K.-T. Park, S. Pilz, S.-J. Seo, A. Voß and J. Zeisig for scientific support, technical assistance and helpful discussions. This work originated within the Joint Research Laboratory of the Korea Institute of Industrial Technology and the Leibniz Institute for Solid State and Materials Research Dresden and funded by the Ministry of Trade, Industry and Energy (MOTIE) of the Republic of Korea on International Collaborative Research and Development Program (N032400011). Additional support through the German Science Foundation (DFG) under the Leibniz Program (Grant EC 111/26-1) and the European Research Council under the ERC Advanced Grant INTELHYB (Grant ERC2013-ADG-340025) is gratefully acknowledged.

References [1] H. Fu, Q. Xiao und, Y. Li, A study of the microstructure and properties of Fe-VW-Mo alloy modified by rare earth, Mater. Sci. Eng. A 395 (2005) 281–287. [2] U. Kühn, N. Mattern, T. Gemming, U. Siegel, K. Werniewicz, J. Eckert, Superior mechanical properties of FeCrMoVC, Appl. Phys. Lett. 90 (2007), 261901/1-3. [3] J. Hufenbach, L. Giebeler, M. Hoffmann, S. Kohlar, U. Kühn, T. Gemming, S. Oswald, B. Eigenmann, J. Eckert, Effect of short-term tempering on microstructure and mechanical properties of high-strength FeCrMoVC, Acta Mater. 60 (2012) 4468–4476. [4] A.S. Chaus, F. Rudnitskii, Structure and properties of cast rapidly cooled highspeed steel R6M5, Met. Sci. Heat Treat. 45 (2003) 157–162. [5] G. Roberts, G. Krauss, R. Kennedy, Tool Steels, fifth ed., ASM International, Materials Park, Ohio, 1998. [6] J. Hufenbach, S. Kohlar, U. Kühn, L. Giebeler, J. Eckert, Microstructural and mechanical characterization of an ultra-high strength Fe86.7Cr4.4Mo0.6V1.1W2.5C4.7 alloy, J. Mater. Sci. 47 (2012) 267–271. [7] M. Bleckmann, J. Gleinig, J. Hufenbach, H. Wendrock, L. Giebeler, J. Zeisig, U. Diekmann, J. Eckert, U. Kühn, Effect of cooling rate on the microstructure and properties of FeCrVC, J. Alloy. Compd. 634 (2015) 200–207. [8] M. Boccalini, H. Goldenstein, Solidification of high speed steels, Int. Mater. Rev. 46 (2001) 92–107. [9] R. Kiessling, N. Lange, Non-Metallic Inclusions in Steel, second ed., Metals

Society, London, 1978. [10] L. Luyckx, The rare earth metals in steel, in: Karl A. Gschneidner Jr. (Ed.), Industrial Applications of Rare Earth Elements, 164, American Chemical Society, Washington, D.C., 1981, pp. 43–78. [11] M. Wang, S. Mu, F. Sun und, Y. Wang, Influence of rare earth elements on microstructure and mechanical properties of cast high-speed steel rolls, J. Rare Earth 25 (2007) 490–494. [12] J. Gao, P. Fu, H. Liu, D. Li, Effects of rare earth on the microstructure and impact toughness of H13 steel, Metals 5 (2015) 383–394. [13] G. Voronin, Thermodynamic characterization of high-temperature superconductors in the yttrium-barium-copper-oxygen system. The Y123 solid solution, Pure Appl. Chem. 72 (2000) 463–477. [14] J. Smialek, B. Pint, Optimizing scale adhesion on single crystal superalloys, Mater. Sci. Forum 369–372 (2001) 459–466. [15] A. Vahed, D. Key, Thermodynamics of rare earths in steelmaking, Metall. Mater. Trans. B 7 (1976) 375–383. [16] H. Fu, Q. Xiao, J. Kuang, Z. Jiang, J. Xing, Effect of rare earth and titanium additions on the microstructures and properties of low carbon Fe-B cast steel, Mater. Sci. Eng. A 466 (2007) 160–165. [17] A.S. Chaus, Use of REM-based modifying agents for improving the structure and properties of cast tungsten-molybdenum high-speed steels, Met. Sci. Heat Treat. 46 (2004) 415–422. [18] P. Waubdy, Rare earth addition to steel, Int. Met. Rev. 1 (1978) 74–98. [19] W. Garrison Jr., J. Maloney, Lanthanum additions and the toughness of ultrahigh strength steels and the determination of appropriate lanthanum additions, Mater. Sci. Eng. A 403 (2005) 299–310. [20] J. Eckert, U. Siegel, U. Kühn, J. Hufenbach, M.-H. Lee, Hochfeste, bei Raumtemperatur plastisch verformbare und mechanische Energie absorbierende Formkörper aus Eisenlegierungen, Patent DE 102010041366 A1. [21] H.M. Rietveld, A profile refinement method for nuclear and magnetic structures, J. Appl. Crystallogr. 2 (1969) 65–71. [22] T. Roisnel, J. Rodríguez-Carvajal, WinPLOTR: a windows tool for powder diffraction patterns analysis, Mater. Sci. Forum 378–381 (2001) 118–123. [23] J. Yang, D. Zou, X. Li, Z. Du, Effect of rare earth on microstructures and properties of high speed steel with high carbon content, J. Iron Steel Res. 14 (2007) 47–52. [24] J.M. Léger, Chalcogenides and pnictides of cerium and uranium under high pressure, Phys. B: Condens. Matter 190B (1993) 84–91. [25] M.L. Gupta, S. Singh, Thermal expansion of CeO2, Ho2O3, and Lu2O3 from 100° to 500° K by an X-ray method, J. Am. Ceram. Soc. 53 (1970) 663–665. [26] H. Bärnighausen, G. Schiller, The crystal structure of A-Ce2O3, J. Less-Common Met. 110 (1985) 385–390. [27] S.P. Ray, D.E. Cox, Neutron diffraction determination of the crystal structure of Ce7O12, J. Solid State Chem. 15 (1975) 333–343. [28] P. Stecher, F. Benesovsky, A. Neckel, H. Nowotny, Untersuchungen in den systemen titan (Zirkonium, Hafnium)-niob-kohlenstoff, Monatshefte für Chemie und verwandte Teile anderer Wissenschaften 12, 1964, pp. 181–195. [29] M. Atoji, Neutron diffraction studies of CaC2, YC2, LaC2, CeC2, TbC2, YbC2, LuC2, and UC2, J. Chem. Phys. 35 (1961) 1950–1960. [30] M. Atoji, D.E. Williams, Neutron-diffraction studies of La2C3, Ce2C3, Pr2C3, and Tb2C3, J. Chem. Phys. 35 (1961) 1960–1966. [31] R. Kohlhaas, P. Dünner, N. Schmitz-Pranghe, Über die temperaturabhängigkeit der gitterparameter von eisen, kobalt und nickel im bereich hoher temperaturen, Z. Angew. Phys. (1967) 245–249. [32] N. Ridley, H. Stuart, Partial molar volumes from high-temperature lattice parameters of iron–carbon austenites, Met. Sci. J. (1970) 219–222. [33] J. Pflüger, J. Fink, W. Weber, K.-P. Bohnen, G. Crecelius, Dielectric properties of TiCx, TiNx, VCx, and VNx from 1.5 to 40 eV determined by electron-energyloss spectroscopy, Phys. Rev. B 30 (1984) 1155–1163. [34] K. Page, J. Li, R. Savinelli, H. Szumila, J. Zhang, J. Stalick, T. Proffen, S. Scott, R. Seshadri, Reciprocal-space and real-space neutron investigation of nanostructured Mo2C and WC, Solid State Sci. 10 (2008) 1499–1510. [35] M. von Stackelberg, Carbide I. Krystallstruktur der carbide MC2, Z. Phys. Chem. B 9 (1930) 437–475. [36] K. Hwang, S. Lee, H. Lee, Effect of alloying elements on microstructure and fracture properties of cast high speed steel rolls, Part II. Fracture behaviour, Mater. Sci. Eng. A 254 (1998) 296–304. [37] I. Picasa, N. Cuadrado, D. Casellas, A. Goez, L. Llanes, Microstructural effects on the fatigue crack nucleation in cold work tool steels, Procedia Eng. 2 (2010) 1777–1785. [38] J. Hirth, M. Cohen, On the strength-differential phenomenon in hardened steel, Metall. Trans. 1 (1970) 3–8.