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Effect of cooling rate on the microstructure and properties of FeCrVC
Journal of Alloys and Compounds 634 (2015) 200–207
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Effect of cooling rate on the microstructure and properties of FeCrVC M. Bleckmann a,⇑, J. Gleinig b, J. Hufenbach b, H. Wendrock b, L. Giebeler b,c, J. Zeisig b, U. Diekmann d, J. Eckert b,c, U. Kühn b a
WIWeB, Institutsweg 1, D-85435 Erding, Germany IFW Dresden, Institute for Complex Materials, P.O. Box 27 01 16, D-01171 Dresden, Germany c TU Dresden, Institute of Materials Science, Helmholtzstraße 7, D-01069 Dresden, Germany d Metatech GmbH, Lünener Straße 211/212, D-59174 Kamen, Germany b
a r t i c l e
i n f o
Article history: Received 24 November 2014 Received in revised form 30 January 2015 Accepted 1 February 2015 Available online 7 February 2015 Keywords: Metals and alloys Thermal analysis Microstructure Phase transitions Thermodynamic modeling
a b s t r a c t In this work a systematic investigation of the influence of the cooling rate on the microstructure and properties of a newly developed Fe92.7Cr4.2V2.1C1 (FeCrVC) tool steel is presented. By applying a tailored casting process and sufficiently high cooling rates excellent mechanical properties are obtained for the presented alloy already in the as-cast state. Since no subsequent heat treatment is required, the cooling parameters applied during the casting process play a key role with respect to the evolving microstructure and resulting properties. In the present publication the effect of the cooling rate on the microstructure and properties of as-solidified FeCrVC was investigated. By using differential scanning calorimetry (DSC), several samples were heated up and cooled with continuous rates of 3–50 K/min. The received DSC data was used to investigate the alloy’s solidification and phase transformation behavior. Subsequently, these samples were studied regarding their properties and microstructure by different analysis methods (EDX/WDX, EBSD, XRD). With increasing cooling rates the liquidus and solidus temperature are lowered, whereas the solidification interval is enlarged. A higher cooling rate is accompanied by a lower solidification time which results in a refinement of the dendritic microstructure. Furthermore, with rising cooling rates the microhardness increased. This provides the opportunity to make predictions from the applied cooling parameters upon the hardness and vice versa and enables one to draw first conclusions on the mechanical properties of the FeCrVC alloy. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Tool steels are widely used and have fundamental relevance in nearly all industrial processes due to their high hardness, strength, and wear resistance [1]. Consequently, there is a strong demand for the development of tool steels with enhanced properties (e.g. by alloy modification or a subsequent heat treatment) and for increasing the efficiency of the manufacturing processes. A promising method for reducing the multi-stage, timeconsuming and cost-intensive manufacturing process of tool steels is near-net-shape casting of the desired components [2–4]. By using a suitable alloy, an additional heat treatment process implying annealing and tempering can be avoided, which is required for many conventional alloys [1]. With this manufacturing technology the microstructure undergoes no further changes after the casting process. This implies that the final properties of the steel result from the microstructure generated during the casting process. ⇑ Corresponding author. Tel.: +49 8122 9590 3260; fax: +49 8122 9590 3202. E-mail address: [email protected] (M. Bleckmann). http://dx.doi.org/10.1016/j.jallcom.2015.02.004 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
As-cast, high alloyed tool steels often exhibit a network of hard eutectic carbides. This carbide network may provide a high hardness and stiffness as well as a better wear resistance compared to wrought tool steels resulting in enhanced tool life [3]. A carbide network may lead to an increased brittleness, especially under tensile load [5]. Beside the chemical composition, the solidification and cooling processes have a considerable effect on the generated microstructure, the resulting properties and the quality of as-cast steels [6–8]. Particularly, the aspect of the cooling rate is crucial as it significantly influences diffusion processes in the liquid and solid state of alloys and was taken into account in former studies [9– 12]. Liquidus and solidus temperatures of steels are not constant, but vary with the cooling rate. Furthermore, an effect on the solidification interval and time is known [13]. Consequently, the morphology, size, and composition of the occurring phases are determined by these two parameters for a given alloy. This mainly affects the dendrite size and distribution of the eutectic [14]. Hereby the secondary dendrite arm spacing k2 can be correlated with the total solidification time tf or cooling rate T_ during solidification,
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which can be described simplified by a power law [9,11,15]. Furthermore, the cooling rate significantly affects the kinetics of phase transformations in solid and controls the developing phases. In many steels low cooling rates lead to equilibrium phases like ferrite, which is soft. In contrast, higher cooling rates promote the non-diffusional martensitic transformation, which is usually preferred due to the occurring hard martensite. For a given alloy, the hardness is generally increased with rising continuous cooling rates [1,16,17]. Therefore, it is most important to properly control solidification and cooling processes of the casting in order to adjust the desired final properties of the material. An alloy qualifying for the near-net-shape casting process, which does not require a subsequent heat treatment is Fe92.7Cr4.2V2.1C1 (FeCrVC). The alloy shows similar mechanical properties like high-alloyed tool steels [18], but without expensive and rare metals like Mo and W as used in many conventional tool steels. Compared to conventional cast steels, the production of this alloy is performed under special conditions including pure elements, inert atmosphere and relatively high cooling rates. Previous investigations demonstrated extraordinary mechanical properties already in the as-cast state of that material, e.g. an engineering compression strength of about 5000 MPa at a fracture strain of 18% and a hardness of 61 HRC [18]. These outstanding characteristics are a result of the special microstructure formed during solidification. It is composed of high strength martensite as well as retained austenite, which is metastable and can transform into martensite during plastic deformation (known as transformation induced plasticity [19]). Furthermore, the present network of complex carbides providing high wear resistance and stiffness predestine the FeCrVC alloy for various applications in tooling [18–20]. The aim of this study was to determine the effect of the cooling rate on the microstructure and properties on as-solidified FeCrVC. Therefore, samples exposed to different cooling rates were examined by several characterization methods. Thereby the influence of the cooling rate on the solidification process, evolving microstructure and resulting microhardness was analyzed.
2. Materials and methods The FeCrVC alloy was prepared by melting pure elements in a ceramic crucible under argon atmosphere in an induction furnace (Balzers). After reaching the casting temperature of 1550 °C, the melt was cast into a copper mold, so that an average solidification rate of 10 to 70 K/s was reached for the resulting ingot (70 mm 120 mm 14 mm) [18]. The actual composition of the ingots was verified by chemical analysis. The carbon content was measured by carrier gas hot extraction (EMIA 820V, Horiba). For the determination of the element contents of iron, chromium, and vanadium inductively coupled plasma optical emission spectrometry (IRIS Intrepid II XUV, Thermo Fisher Scientific) was used. Table 1 displays the nominal chemical composition of the FeCrVC alloy as well as the average content of Fe, Cr, V, and C resulting from a threefold determination. Thereby the nominal composition of the alloy and the real element contents demonstrate a good correlation. Out of the ingot several samples were cut for thermal analysis (DSC), which was performed using a Netzsch 404C. To exclude a pronounced mass influence, the weight of the samples was restricted to approx. 20 mg. In dynamic measurements, these samples were heated above the liquidus temperature to 1480 °C, held for 10 min and subsequently cooled to room temperature. The heating and cooling processes were carried out under argon flow (50 ml/min), using scanning rates of 3, 5, 10, 20, 30, 40, and 50 K/min. To check the reproducibility of the DSC signal every scanning rate was measured for two samples. Prior to the measurements the DSC
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was calibrated for each scanning rate using pure iron, gold, and indium. The system was additionally referenced with two empty corundum crucibles for each scanning rate. After DSC measurements the samples were metallographically prepared and etched with Nital solution (6.5 ml HNO3 and 93.5 ml C2H5OH). Afterward, the microstructure of the samples was examined by optical microscopy (OM; Epiphot 300, Nikon). The measurement of the secondary dendrite arm spacings (SDAS) was performed according to Ref. [21] as far as possible. At higher magnifications, scanning electron microscopy was used (SEM; 1530 Gemini, Zeiss) in combination with energy dispersive X-ray spectroscopy (EDX; XMAX80, Oxford Instruments) for detecting the element distribution. For the identification of the crystal structure of carbides, electron backscatter diffraction (EBSD; Channel 5, Nordlys 2, 20 kV, tilt 70°) was performed. For cross-checking the EBSD results and additional structural characterization, X-ray diffraction (XRD; Stadi P, STOE, Dectris Mythen 1 K detector) with the second charge of DSC samples was performed in transmission geometry using Mo Ka1 radiation. The preparation of the samples for XRD and EBSD is described in Ref. [18]. Vickers microhardness HV0.1 of the first batch of DSC samples was measured using a hardness tester (HMV-2000, Shimadzu). JMatPro, a commercial software based on CalPhaD, was used for the thermodynamic calculation of the FeCrVC phase diagram [22].
3. Results and discussion 3.1. Thermodynamic calculation To predict the phases occurring in the FeCrVC alloy during the solidification and cooling process, a thermodynamic calculation using JMatPro was performed. In the temperature range between 20 and 1600 °C the phase diagram (equilibria) for the FeCrVC alloy was calculated from the nominal chemical composition given in Table 1. The result displayed in Fig. 1 shows the obtained mass fraction of phases related to the temperature. At 1445 °C the solidification of FeCrVC starts with the formation of an austenitic phase. Around 1290 °C the solidification is completed and simultaneously the precipitation of MC carbides begins. With decreasing temperature at 940 °C a second type of carbide with the constitution M7C3 starts to precipitate. Under equilibrium cooling conditions the remaining austenite transforms between 789 and 767 °C entirely into a-ferrite, exhibiting a bcc structure. Further cooling, down to 20 °C, proceeds without further transformations. At room temperature and under equilibrium conditions the final microstructure would consist of 91 wt% ferrite, 6.5 wt% M7C3, and 2.5 wt% MC carbides. The calculated composition of the phases at room temperature is displayed in Table 2. Ferrite consists basically of iron and a small amount of chromium, while it contains almost no carbon and vanadium. Beside carbon and chromium, the MC carbides consists mainly vanadium but no iron. The M7C3 carbides are based on chromium and carbon, as well as a substantial amount of iron. Altogether the calculated
Table 1 Nominal and real chemical composition of the investigated FeCrVC alloy. Element
Fe
Cr
V
C
Nominal (wt%) Real (wt%)
92.7 92.66
4.2 4.16
2.1 2.06
1.0 0.99
Fig. 1. Mass fraction of equilibrium phases of the FeCrVC-alloy as a function of temperature, simulated with JMatPro.
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Table 2 Simulated constitution of FeCrVC phases in equilibrium at room temperature. Content (wt%)
Fe
Cr
V
C
Ferrite MC carbide M7C3 carbide
99.6 0.0 33.2
0.4 4.6 56.1
0.1 77.4 1.9
0.1 18.0 8.8
carbide types and composition for an Fe–Cr–V–C alloy coincide with the literature in Refs. [23–25]. 3.2. Microstructure analysis After considering the theoretical phase diagram of FeCrVC with help of the simulation in equilibrium, the actual circumstances were observed by a close examination of the DSC samples. In order to evaluate the effect of the cooling rate on the solidification morphology and evolving microstructure, several specimens were linearly cooled down from the liquid state with different rates (3–50 K/min) from 1480 °C. Fig. 2 shows a comparison of the ascast microstructures of the DSC samples cooled at different rates. The specimen with a rate of 3 K/min (Fig. 2a) cooled down very slowly and shows a globular morphology. According to the XRD pattern presented in Fig. 3a the matrix phase exhibits a bcc structure with a lattice parameter a = 0.287081(6) nm. Hereby the structure model required for the refinement was taken from Ref. [18]. Together with the low matrix microhardness (about 188 HV0.1), the cubic structure can be presumed as ferrite [23]. Based on the globular ferritic microstructure this DSC sample shows parallels with the simulation (Fig. 1), which indicates an austenite into ferrite transformation at a temperature of nearly 800 °C for solidification and cooling rates near equilibrium. Furthermore, in microscopic images a small amount of carbides at the grain boundaries forming a partly connected network is recognizable (see inset of Fig. 2a). They are rich in V and were assumed as MC carbides with M = V and Cr (not shown here) but were not detectable in the XRD pattern. Due to the carbide morphology, their position at the grain boundaries as well as the results of previous studies, the carbides can be identified as MC type [18,19,26,27]. M7C3 carbides like displayed in the simulation (Fig. 1) were not found by applied techniques. In contrast, alloy specimens cooled down with 10–50 K/min (Fig. 2b–f) exhibit a dendritic structure. On the basis of the XRD pattern, exemplarily shown for the 20 K/min sample in Fig. 3b, the included crystal structures of the dendritic DSC specimens were identified. For the determination of their lattice parameters the present phases were indexed based on structure models according to Ref. [18]. The patterns predominant peaks represent a bcc crystal structure exhibiting a lattice parameter a = 0.28720(3) nm, which could refer to martensite. The significantly broadened XRD peaks compared to the ferritic specimen and slightly larger lattice parameter give hints for a lattice distortion and slight expansion due to the carbon dissolved in martensite [19]. Furthermore, the detected fcc structure can be assigned to an austenitic phase with a lattice parameter of a = 0.3601(4) nm. Beside these two main phases, there are signs in the XRD pattern indicating the existence of carbides. Due to their low content the carbides are at the detection limit, only their main peak is visible fairly poor, so they were neglected in the performed Rietveld analysis. As a result, the sample consists of 92 wt% martensite and 8 wt% austenite. Fig. 4 shows exemplarily the microstructure of the dendritic DSC sample cooled down with 20 K/min at higher magnifications. Hereby the dark brown structural constituents in the dendritic area are martensite plates. They are embedded in a bright appearing phase, which is composed of retained austenite. Furthermore a
small volume of carbides, visible as skeleton-like light gray phases, is situated at the former austenite grain boundaries. The carbides are clearly visible as dark phase in the SEM image in backscattered electrons (BSE)-contrast in Fig. 5a. Thereby two types of carbides with varying shapes and gray scales were observed indicating different chemical compositions. As their volume fraction was too low for a clear identification in the XRD pattern, the characterization of the carbides was done by EDX and EBSD. Based on the EDX results shown in Fig. 5b–e, two types of carbides can be distinguished: those that are rich in V and those that exhibit mainly Cr with notable amount of Fe. Typical EBSD diffraction patterns taken from the V-rich and Cr-rich carbides are displayed in Fig. 6. The experimental results are in accordance with the found indexed patterns showing the Miller indices of the zone axes. The V-rich carbides accord well with a fcc crystal structure referring to the MC-type. For the Cr-rich carbides a good accordance with the orthorhombic M7C3 structure was detected. The analyzed carbide types are in agreement with the results of the simulation. Apart from the carbides, single martensite plates are recognizable in the SEM image (Fig. 5a). In addition, the enrichment of the austenite with chromium can be seen in the EDX data (Fig. 5b–e) as reported in Ref. [19]. By means of the previous examinations the influence of the cooling rate on the evolving phases of FeCrVC is shown. Only slow cooling rates lead to a ferritic matrix according to the simulation displayed in Fig. 1, whereas at rates higher than 10 K/min a microstructure mainly composed of martensite and retained austenite is already formed. Because the cooling rate determines the evaluation of phases in the alloy, it is anticipated that the mechanical properties will likewise be influenced by cooling rate. This is confirmed by the results of the microhardness measurement in Table 3. A significant rise in the microhardness of the DSC samples between the samples cooled with 5 K/min from about 180 HV0.1 to more than 450 HV0.1 is observed. This is attributed to the change of the main phase from soft ferrite to harder martensite [23]. Thereafter DSC samples with proceeding increase of the cooling rate exhibit a continuously rising microhardness. It is assumed to be caused by an increasing strength of the martensitic phase. The reason for this could be increasing deformations in the martensite lattice due to higher contents of carbon and alloying elements [17]. This can be concluded from the simultaneous drop of the martensite transformation temperature, which is described below (see Section 3.3.). Beside the effect of the cooling rate on the occurring phases and resulting hardness, an impact on the morphology size is observed, too. The dendritic structure (Fig. 2) is continuously refined with increasing cooling rate, which was quantified by measuring the SDAS. Due to the small specimen size in relation to the size of the dendritic structure each DSC sample contained only few dendrites. Hence less than five SDAS could be captured whereupon the requirements for an accurate measurement according to Ref. [21] were fulfilled only partially. The result of the measured SDAS as arithmetic mean were plotted against the corresponding cooling rate during solidification in a double-logarithmic scale. Hereby SDAS k2 and cooling rate T_ were found to be related by the equation
k2 ¼ AT_ b : The parameters were determined as A = 285.9 lm and b = 0.398 for cooling rate values in K/min. The rising cooling rate is accompanied by a smaller solidification time. As the diffusion in the melt is higher than in the solid, this is followed by depression of the Ostwald ripening process eventually leading to smaller dendrites [28].
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Fig. 2. OM images of as-solidified FeCrVC-samples cooled with 3 (a), 10 (b), 20 (c), 30 (d), 40 (e) and 50 K/min (f), etched with Nital; inset in (a) shows SEM image of carbides along the grain boundary in FeCrVC-sample.
Fig. 3. Experimental XRD pattern of the DSC samples with cooling rates of 3 K/min (a) and 20 K/min (b) with identified phases, calculated fit (Rietveld method) and difference plot (difference of calculated and experimental curve).
3.3. Differential scanning calorimetry Prior to the cooling process the as-cast specimens were heated into the liquid state up to 1480 °C. The obtained DSC heating and cooling curves show several endothermic and exothermic peaks,
which are attributed to phase transformations in the specimen. The melting process is described by the endothermic peaks of the heating curve, whereas the cooling curve exhibits exothermic peaks. Fig. 7a shows a comparison of the heat flow signals for different cooling curves. The two DSC curves obtained for each
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Fig. 4. OM image of the as-solidified structure taken from the 20 K/min sample showing the dendritic morphology composed of martensite, austenite and carbides, etched with Nital.
cooling rate were found to be consistent, certifying the measured reaction sequences as reliable. Therefore, only one representative curve for each cooling rate was displayed in Fig. 7a. Due to the focus on the solidification behavior, a typical DSC cooling curve (cooling rate 20 K/min) is displayed in detail in Fig. 7b and c. The solidification part of the DSC curve in Fig. 7b shows two major peaks, whereas the first one is a double peak. The start of the solidification is characterized by the onset of the
Fig. 5. SEM image of the 20 K/min sample taken in BSE-contrast (a) with related EDX maps of iron (b), chromium (c), vanadium (d) and carbon (e).
first event, which represents the liquidus temperature Tl. It can be assumed that this first peak section can be attributed to the formation of primary delta ferrite. Although the simulation for the alloy composition (Fig. 1) does not predict a formation of delta ferrite, the typical delta ferrite peak [11] is visible for slow heating and cooling rates up to 20 K/min (Fig. 7a). This can be mainly attributed due to carbon escaping into the atmosphere during the DSC measurements. At slow heating and cooling rates, a larger quantity of carbon escapes, so that the phase formation calculated for an alloy with 1 wt% of carbon (Fig. 1) is not valid anymore and the formation of delta ferrite becomes possible. For faster DSC measurements (P30 K/min) less carbon can evaporate. This leads to a direct transformation of liquid into austenite and, therefore, no double peak is formed (as shown in Fig. 7b). To verify this assumption the final carbon content of several DSC samples was quantified by carrier gas hot extraction directly after the performed DSC measurements. For heating and cooling rates of 20 K/min the specimens contain in average 0.48 wt% carbon and for 40 K/min an average carbon content of 0.8 wt% was determined. These measurements show that the final carbon content is significantly influenced by the heating and cooling rate of the DSC sample and, thus, the amount of carbon escaping during the DSC measurements essentially affects the solidification path of the alloy. Fig. 8 shows the mass fraction of the equilibrium phases as a function of the temperature for a similar alloy where the carbon content is 0.5 wt% and the iron-content is slightly higher (Fe93.2Cr4.2V2.1C0.5) compared to the base alloy. It is shown that delta ferrite is formed, which is consistent with the results of the DSC measurements. It can be assumed that cooling rates slower than 20 K/min do have the same or even less carbon content. In comparison to microscopic results of the DSC sample cooled down with 3 K/ min, the calculations with 0.5 wt% C (Fig. 8) show more reliable formation pathes and phase fraction values for the arising carbides than the simulation for 1 wt% C (Fig. 1). After solidification and fully formation of an austenitic matrix carbides of the MC type form at the austenitic grain boundaries as shown in Fig. 2a. M7C3 carbides are generated relatively late in the solid state after the transformation of austenite into ferrite. The simulation with 0.5 wt% C (Fig. 8) depicts for the M7C3 carbides a phase fraction which is distinctly less than in the first calculation (Fig. 1). These carbides should be very fine and hardly detectable by XRD or microscopic measurements which was already assumed in Section 3.2. The second part of the double peak in Fig. 7b embodies the peritectic transformation of liquid and delta ferrite in austenite [11,25]. Hereafter, the second peak represents an eutectic reaction where the remaining melt is decomposed into austenite and MxC carbides [11,29], which have a MC structure as shown by microstructure analysis where the eutectic MC (V rich) carbides are embedded in austenitic phase along the grain boundaries (Fig. 4). In this study the offset of the second peak is taken as solidus temperature Ts. The further DSC cooling curve in the solid state is displayed in Fig. 7c. There are three exothermic peaks, which exhibit a much lower heat flow than the solidification peaks. It can be assumed that the peak at 1100 °C represents the precipitation of carbides with a M7C3 (Cr rich) structure found in the DSC sample. After the solidification was completed this carbide type could form out of the Cr enriched matrix on the energetically favoured interfaces of eutectic MC carbides in a very fine morphology as shown in Fig. 5a. A further minor peak arises around 700 °C, which is typical for the austenite-ferrite transformation. It is possible that a small amount of bainite structure has formed, which has to be checked by further examinations. The large peak between 300 and 400 °C can be attributed to the formation of martensitic phase, which is
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Fig. 6. Experimental EBSD patterns taken from the vanadium (a) and chromium rich carbides (c) fit with the indexed patterns of fcc-type VC (b) and orthorhombic Cr7C3 (d).
Table 3 Microhardness of the DSC sample cooled with 5 K/min measured in ferrite grains as well as measured microhardness in the dendritic area of the alloy mainly composed of martensite and retained austenite (10–50 K/min). Cooling rate (K/min)
5
10
20
30
40
50
Microhardness (HV0.1)
173 ± 40
456 ± 19
498 ± 13
594 ± 18
651 ± 28
748 ± 31
Fig. 7. DSC cooling curves of FeCrVC at various scanning rates in comparison (a) showing the position of liquidus Tl, solidus Ts and martensite start temperature Ms; DSC curve for a cooling rate of 20 K/min showing the solidification (b) and cooling in the solid state (c).
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Fig. 8. Mass fraction of equilibrium phases of Fe92.2Cr4.2V2.1C0.5 as a function of temperature, simulated with JMatPro.
contained in the 20 K/min DSC sample (Fig. 4). Hereby the peak onset is characterized as martensite start temperature Ms. In general the heat flow signal increases with higher cooling rates due to the rising enthalpy [8,12]. For this reason signals from faster cooling rates have a better visibility but also tend to overlap and conceal single reactions [30]. Thus the peaks for the delta ferritic and austenitic solidification can be clearly distinguished up to 20 K/min, while at higher rates only one major peak is visible. Apart from the solidification peak, signals with cooling rates of 3 and 5 K/min exhibit around 800 °C another evanescent peak, which is visible only in detail (not shown here). The DSC samples cooled with 3 and 5 K/min have a very low microhardness due to the formation of the ferritic phase. Hence the peak in the heat flow can be attributed to the transformation from austenite into ferrite, which was proposed by the simulation in the same range of temperatures. DSC curves with cooling rates of 10 K/min and above features another peak arising around 400 °C (Fig. 7a). This peak can be attributed to the formation of martensite, which is present in the mentioned alloy specimens. By examining the identical peaks at different cooling rates a general shift toward lower temperatures at higher cooling rates is observed. As a consequence, the characteristic temperatures Tl, Ts, and Ms varying with the cooling rate. Thereby the deviation of the martensite start temperature can be mainly assigned to the escaping carbon during the DSC measurements at slower heating and cooling rates. The results in Fig. 9 illustrates that an increasing cooling rate leads to lower liquidus and solidus temperatures. DSC heating curves do not show this dependency. Hereby the dependence between the liquidus and respectively solidus temperature and DSC cooling rate can be described as linear approximation. By linear extrapolation it was possible to estimate liquidus and solidus temperature at a zero velocity in equilibrium [13]. The values found are almost equal for the heating and cooling process as seen in Table 4. Furthermore, the experimentally determined temperatures were found to be in good agreement with values of the simulation and approximated calculations. Thereby the approximated liquidus temperature was determined by the empirical formula given in Ref. [31] via the chemical constitution of FeCrVC:
Fig. 9. Influence of the cooling and heating rate on the liquidus (a) and solidus temperature (b) of FeCrVC.
Table 4 Liquidus and solidus temperature of FeCrVC in equilibrium, linear extrapolation of data, simulation by JMatPro, approximation after formula in Ref. [31].
Tl (°C) Ts (°C)
DSC-heating
DSC-cooling
Simulation
Approximation
1456 1283
1453 1284
1445 1290
1455 –
However, the solidification time tf obviously decreases at higher cooling rates due to the accelerated cooling process. Hereby solidification time and cooling rate are related by an inverse proportionality. 4. Summary and outlook The present work demonstrates that the cooling rate has a distinctive influence on the microstructure and properties of the investigated alloy Fe92.7Cr4.2V2.1C1.
T l ð CÞ ¼ 1531—61:5C—4Mn—14Si—45S—30P—1:5Cr —2:5Al—3:5Ni—4V—5Mo ðcontents in wt%Þ: Altogether the observed drop of solidus temperature is larger than for the liquidus temperature. This leads to a higher solidification interval DT with increase of the cooling rate, which was found to rise linearly as presented in Fig. 10.
Fig. 10. Solidification time tf and interval DT of FeCrVC as a function of cooling rate.
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In this study, DSC samples were heated and solidified at various cooling rates between 3 and 50 K/min. With the help of the obtained alloy specimens and appropriate DSC signal curves, the influence of the cooling rate was systematically analyzed and characterized. The results show that the temperature control during DSC measurements significantly affects the solidification process. Liquidus and solidus temperature are lowered with increasing cooling rates whereas the solidification interval, which is important for the processing of the material, is expanding. The estimated equilibrium liquidus and solidus temperatures are 1453 °C and 1284 °C respectively. Furthermore, an influence on the evolving phases and morphology was observed. Specimens manufactured at slow cooling rates solidify in a cellular structure and consist mainly of ferrite. In contrast, specimens cooled at higher rates exhibiting dendritic microstructure composed of the non-equilibrium phases martensite and retained austenite. Simultaneously a refinement of the dendritic morphology as well as a depression of the martensite start temperature with an increasing cooling rate was observed. Beside the main phases, a network of mixed carbides was detected and verified as MC (M = V, Cr) and M7C3 (M = Cr, Fe). Moreover, the investigated alloy exhibits a rising microhardness at increasing cooling rates. Altogether, the presented work clarifies the connection between cooling rate, secondary dendrites arm spacing and resulting microhardness for a newly developed tool steel. This allows to draw first conclusions from the solidification parameters with regard to the resulting mechanical properties. Thereby it could be displayed that the carbon content is significantly influenced by the temperature regime and therefore affects the phase transformations. The gained knowledge is a step toward systematically adjusting and enhancing the properties of as-solidified tool steels by tailored temperature control. The equilibrium phase simulation consulted for comparison provided good results regarding the solidification parameters and occurring phases. In future the simulation results will be combined with the displayed parameter relations. This should provide the opportunity for more reliable simulations in the non-equilibrium state and offers novel possibilities for predictions close to real production processes. Acknowledgement The authors would like to thank B. Bartusch, H. Bußkamp, S. Donath, M. Frey, C. Geringswald, W. Gruner, V. Hammond, R. Keller,
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and A. Voß for scientific support, technical assistance and helpful discussions. References [1] G. Roberts, G. Krauss, R. Kennedy, Tool Steels, fifth ed., ASM International, 1998. [2] A.S. Chaus, F.I. Rudnitskii, Met. Sci. Heat Treat. 45 (2003) 157–162. [3] A.S. Chaus, ISIJ Int. 45 (2005) 1297–1306. [4] W. Löser, S. Thiem, M. Jurisch, Mater. Sci. Eng. A 173 (1993) 323–326. [5] A. Akbarzadeh, S. Naghdy, Mater. Des. 46 (2013) 654–659. [6] K. Gryc, B. Smetana, M. Zˇaludová, K. Michalek, M. Tkadlecˇková, J. Dobrovská, L. Socha, L. Válek, R. Pachlopnik, Metalurgija 52 (2013) 445–448. [7] D. Steiner Petrovicˇ, M. Pirnat, G. Klancˇnik, P. Mrvar, J. Medved, J. Therm. Anal. Calorim. 109 (2012) 1185–1191. [8] D. Steiner Petrovicˇ, G. Klancˇnik, M. Pirnat, J. Medved, J. Therm. Anal. Calorim. 105 (2011) 251–257. [9] S. Gowri, F.H. Samuel, Metall. Trans. A 23 (1992) 3369–3376. [10] S. Milenkovic, I. Sabirov, J. Llorca, Mater. Lett. 73 (2012) 216–219. [11] H.F. Fischmeister, R. Riedl, S. Karagoz, Metall. Trans. A 20 (1989) 2133–2148. [12] M.N. Khan, M. Aljarrah, J.T. Wood, M. Medraj, J. Mater. Res. 26 (2011) 974–982. [13] B. Smetana, M. Zˇaludová, S. Zlá, J. Dobrovská, M. Cagala, I. Szurman, D. Petlák, Application of high temperature data technique to fe based systems, in: Paper presented at. METAL 2010 Conference, Rozˇnov pod Radhošteˇm, Czech Republic, 2010 [14] M. Boccalini, H. Goldenstein, Int. Mater. Rev. 46 (2001) 92–115. [15] M.A. Taha, J. Mater. Sci. Lett. 5 (1986) 307–310. [16] G.M. Vorobev, A.F. Siukhin, V.I. Popova, Met. Sci. Heat Treat. 9 (1967) 325–326. [17] G.F. Vander Voort (Ed.), Atlas of Time-temperature Diagrams for Irons and Steels, ASM International, 2007. [18] J. Hufenbach, K. Kunze, L. Giebeler, T. Gemming, H. Wendrock, C. Baldauf, U. Kühn, W. Hufenbach, J. Eckert, Mater. Sci. Eng. A 586 (2013) 267–275. [19] U. Kühn, J. Romberg, N. Mattern, H. Wendrock, J. Eckert, J. Mater. Res. 25 (2010) 368–374. [20] A.S. Chaus, M. Hudáková, Wear 267 (2009) 1051–1055. [21] Jernkontoret, A Guide to the Solidification of Steels, Ljungberg Tryckeri AB, Stockholm, 1977. [22] U. Diekmann, Calculation of steel data using JMatPro, in: Paper presented at ˇ , Czech Republic, 2012. COMAT 2012 Conference, Plzen [23] M. Durand-Charre, Microstructures of Steels and Cast Irons, Springer, Berlin, 2004. [24] J. Bratberg, K. Frisk, Metall. Mater. Trans. A 35 (2004) 3649–3663. [25] V. Raghavan, Phase Diagrams of Quaternary Iron Alloys, The Indian Institute of Metals, Kolkata, 1996. [26] V. Vitry, S. Nardone, J.-P. Breyer, M. Sinnaeve, F. Delaunois, Mater. Des. 34 (2012) 372–378. [27] J. Hufenbach, U. Kühn, L. Krüger, H. Wendrock, J. Eckert, Steel Res. Int. 82 (2011) 51–55. [28] W. Kurz, D.J. Fisher, Fundamentals of Solidification, fourth ed., Trans Tech Publications, 1998. [29] D.A. Porter, K.E. Easterling, Phase Transformations in Metals and Alloys, second ed., Nelson Thornes, 2001. [30] L.C. Thomas, Am. Lab. 33 (2001) 26–29. [31] E.T. Turkdogan, R.J. Fruehan, Fundamentals of Iron and Steelmaking, in: R.J. Fruehan (Ed.), The Making, Shaping and Treating of Steels, 11th ed., The AISE Steel Foundation, Pittsburgh, 1998.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Effect of cooling rate on the microstructure and properties of FeCrVC M. Bleckmann a,⇑, J. Gleinig b, J. Hufenbach b, H. Wendrock b, L. Giebeler b,c, J. Zeisig b, U. Diekmann d, J. Eckert b,c, U. Kühn b a
WIWeB, Institutsweg 1, D-85435 Erding, Germany IFW Dresden, Institute for Complex Materials, P.O. Box 27 01 16, D-01171 Dresden, Germany c TU Dresden, Institute of Materials Science, Helmholtzstraße 7, D-01069 Dresden, Germany d Metatech GmbH, Lünener Straße 211/212, D-59174 Kamen, Germany b
a r t i c l e
i n f o
Article history: Received 24 November 2014 Received in revised form 30 January 2015 Accepted 1 February 2015 Available online 7 February 2015 Keywords: Metals and alloys Thermal analysis Microstructure Phase transitions Thermodynamic modeling
a b s t r a c t In this work a systematic investigation of the influence of the cooling rate on the microstructure and properties of a newly developed Fe92.7Cr4.2V2.1C1 (FeCrVC) tool steel is presented. By applying a tailored casting process and sufficiently high cooling rates excellent mechanical properties are obtained for the presented alloy already in the as-cast state. Since no subsequent heat treatment is required, the cooling parameters applied during the casting process play a key role with respect to the evolving microstructure and resulting properties. In the present publication the effect of the cooling rate on the microstructure and properties of as-solidified FeCrVC was investigated. By using differential scanning calorimetry (DSC), several samples were heated up and cooled with continuous rates of 3–50 K/min. The received DSC data was used to investigate the alloy’s solidification and phase transformation behavior. Subsequently, these samples were studied regarding their properties and microstructure by different analysis methods (EDX/WDX, EBSD, XRD). With increasing cooling rates the liquidus and solidus temperature are lowered, whereas the solidification interval is enlarged. A higher cooling rate is accompanied by a lower solidification time which results in a refinement of the dendritic microstructure. Furthermore, with rising cooling rates the microhardness increased. This provides the opportunity to make predictions from the applied cooling parameters upon the hardness and vice versa and enables one to draw first conclusions on the mechanical properties of the FeCrVC alloy. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Tool steels are widely used and have fundamental relevance in nearly all industrial processes due to their high hardness, strength, and wear resistance [1]. Consequently, there is a strong demand for the development of tool steels with enhanced properties (e.g. by alloy modification or a subsequent heat treatment) and for increasing the efficiency of the manufacturing processes. A promising method for reducing the multi-stage, timeconsuming and cost-intensive manufacturing process of tool steels is near-net-shape casting of the desired components [2–4]. By using a suitable alloy, an additional heat treatment process implying annealing and tempering can be avoided, which is required for many conventional alloys [1]. With this manufacturing technology the microstructure undergoes no further changes after the casting process. This implies that the final properties of the steel result from the microstructure generated during the casting process. ⇑ Corresponding author. Tel.: +49 8122 9590 3260; fax: +49 8122 9590 3202. E-mail address: [email protected] (M. Bleckmann). http://dx.doi.org/10.1016/j.jallcom.2015.02.004 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
As-cast, high alloyed tool steels often exhibit a network of hard eutectic carbides. This carbide network may provide a high hardness and stiffness as well as a better wear resistance compared to wrought tool steels resulting in enhanced tool life [3]. A carbide network may lead to an increased brittleness, especially under tensile load [5]. Beside the chemical composition, the solidification and cooling processes have a considerable effect on the generated microstructure, the resulting properties and the quality of as-cast steels [6–8]. Particularly, the aspect of the cooling rate is crucial as it significantly influences diffusion processes in the liquid and solid state of alloys and was taken into account in former studies [9– 12]. Liquidus and solidus temperatures of steels are not constant, but vary with the cooling rate. Furthermore, an effect on the solidification interval and time is known [13]. Consequently, the morphology, size, and composition of the occurring phases are determined by these two parameters for a given alloy. This mainly affects the dendrite size and distribution of the eutectic [14]. Hereby the secondary dendrite arm spacing k2 can be correlated with the total solidification time tf or cooling rate T_ during solidification,
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which can be described simplified by a power law [9,11,15]. Furthermore, the cooling rate significantly affects the kinetics of phase transformations in solid and controls the developing phases. In many steels low cooling rates lead to equilibrium phases like ferrite, which is soft. In contrast, higher cooling rates promote the non-diffusional martensitic transformation, which is usually preferred due to the occurring hard martensite. For a given alloy, the hardness is generally increased with rising continuous cooling rates [1,16,17]. Therefore, it is most important to properly control solidification and cooling processes of the casting in order to adjust the desired final properties of the material. An alloy qualifying for the near-net-shape casting process, which does not require a subsequent heat treatment is Fe92.7Cr4.2V2.1C1 (FeCrVC). The alloy shows similar mechanical properties like high-alloyed tool steels [18], but without expensive and rare metals like Mo and W as used in many conventional tool steels. Compared to conventional cast steels, the production of this alloy is performed under special conditions including pure elements, inert atmosphere and relatively high cooling rates. Previous investigations demonstrated extraordinary mechanical properties already in the as-cast state of that material, e.g. an engineering compression strength of about 5000 MPa at a fracture strain of 18% and a hardness of 61 HRC [18]. These outstanding characteristics are a result of the special microstructure formed during solidification. It is composed of high strength martensite as well as retained austenite, which is metastable and can transform into martensite during plastic deformation (known as transformation induced plasticity [19]). Furthermore, the present network of complex carbides providing high wear resistance and stiffness predestine the FeCrVC alloy for various applications in tooling [18–20]. The aim of this study was to determine the effect of the cooling rate on the microstructure and properties on as-solidified FeCrVC. Therefore, samples exposed to different cooling rates were examined by several characterization methods. Thereby the influence of the cooling rate on the solidification process, evolving microstructure and resulting microhardness was analyzed.
2. Materials and methods The FeCrVC alloy was prepared by melting pure elements in a ceramic crucible under argon atmosphere in an induction furnace (Balzers). After reaching the casting temperature of 1550 °C, the melt was cast into a copper mold, so that an average solidification rate of 10 to 70 K/s was reached for the resulting ingot (70 mm 120 mm 14 mm) [18]. The actual composition of the ingots was verified by chemical analysis. The carbon content was measured by carrier gas hot extraction (EMIA 820V, Horiba). For the determination of the element contents of iron, chromium, and vanadium inductively coupled plasma optical emission spectrometry (IRIS Intrepid II XUV, Thermo Fisher Scientific) was used. Table 1 displays the nominal chemical composition of the FeCrVC alloy as well as the average content of Fe, Cr, V, and C resulting from a threefold determination. Thereby the nominal composition of the alloy and the real element contents demonstrate a good correlation. Out of the ingot several samples were cut for thermal analysis (DSC), which was performed using a Netzsch 404C. To exclude a pronounced mass influence, the weight of the samples was restricted to approx. 20 mg. In dynamic measurements, these samples were heated above the liquidus temperature to 1480 °C, held for 10 min and subsequently cooled to room temperature. The heating and cooling processes were carried out under argon flow (50 ml/min), using scanning rates of 3, 5, 10, 20, 30, 40, and 50 K/min. To check the reproducibility of the DSC signal every scanning rate was measured for two samples. Prior to the measurements the DSC
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was calibrated for each scanning rate using pure iron, gold, and indium. The system was additionally referenced with two empty corundum crucibles for each scanning rate. After DSC measurements the samples were metallographically prepared and etched with Nital solution (6.5 ml HNO3 and 93.5 ml C2H5OH). Afterward, the microstructure of the samples was examined by optical microscopy (OM; Epiphot 300, Nikon). The measurement of the secondary dendrite arm spacings (SDAS) was performed according to Ref. [21] as far as possible. At higher magnifications, scanning electron microscopy was used (SEM; 1530 Gemini, Zeiss) in combination with energy dispersive X-ray spectroscopy (EDX; XMAX80, Oxford Instruments) for detecting the element distribution. For the identification of the crystal structure of carbides, electron backscatter diffraction (EBSD; Channel 5, Nordlys 2, 20 kV, tilt 70°) was performed. For cross-checking the EBSD results and additional structural characterization, X-ray diffraction (XRD; Stadi P, STOE, Dectris Mythen 1 K detector) with the second charge of DSC samples was performed in transmission geometry using Mo Ka1 radiation. The preparation of the samples for XRD and EBSD is described in Ref. [18]. Vickers microhardness HV0.1 of the first batch of DSC samples was measured using a hardness tester (HMV-2000, Shimadzu). JMatPro, a commercial software based on CalPhaD, was used for the thermodynamic calculation of the FeCrVC phase diagram [22].
3. Results and discussion 3.1. Thermodynamic calculation To predict the phases occurring in the FeCrVC alloy during the solidification and cooling process, a thermodynamic calculation using JMatPro was performed. In the temperature range between 20 and 1600 °C the phase diagram (equilibria) for the FeCrVC alloy was calculated from the nominal chemical composition given in Table 1. The result displayed in Fig. 1 shows the obtained mass fraction of phases related to the temperature. At 1445 °C the solidification of FeCrVC starts with the formation of an austenitic phase. Around 1290 °C the solidification is completed and simultaneously the precipitation of MC carbides begins. With decreasing temperature at 940 °C a second type of carbide with the constitution M7C3 starts to precipitate. Under equilibrium cooling conditions the remaining austenite transforms between 789 and 767 °C entirely into a-ferrite, exhibiting a bcc structure. Further cooling, down to 20 °C, proceeds without further transformations. At room temperature and under equilibrium conditions the final microstructure would consist of 91 wt% ferrite, 6.5 wt% M7C3, and 2.5 wt% MC carbides. The calculated composition of the phases at room temperature is displayed in Table 2. Ferrite consists basically of iron and a small amount of chromium, while it contains almost no carbon and vanadium. Beside carbon and chromium, the MC carbides consists mainly vanadium but no iron. The M7C3 carbides are based on chromium and carbon, as well as a substantial amount of iron. Altogether the calculated
Table 1 Nominal and real chemical composition of the investigated FeCrVC alloy. Element
Fe
Cr
V
C
Nominal (wt%) Real (wt%)
92.7 92.66
4.2 4.16
2.1 2.06
1.0 0.99
Fig. 1. Mass fraction of equilibrium phases of the FeCrVC-alloy as a function of temperature, simulated with JMatPro.
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Table 2 Simulated constitution of FeCrVC phases in equilibrium at room temperature. Content (wt%)
Fe
Cr
V
C
Ferrite MC carbide M7C3 carbide
99.6 0.0 33.2
0.4 4.6 56.1
0.1 77.4 1.9
0.1 18.0 8.8
carbide types and composition for an Fe–Cr–V–C alloy coincide with the literature in Refs. [23–25]. 3.2. Microstructure analysis After considering the theoretical phase diagram of FeCrVC with help of the simulation in equilibrium, the actual circumstances were observed by a close examination of the DSC samples. In order to evaluate the effect of the cooling rate on the solidification morphology and evolving microstructure, several specimens were linearly cooled down from the liquid state with different rates (3–50 K/min) from 1480 °C. Fig. 2 shows a comparison of the ascast microstructures of the DSC samples cooled at different rates. The specimen with a rate of 3 K/min (Fig. 2a) cooled down very slowly and shows a globular morphology. According to the XRD pattern presented in Fig. 3a the matrix phase exhibits a bcc structure with a lattice parameter a = 0.287081(6) nm. Hereby the structure model required for the refinement was taken from Ref. [18]. Together with the low matrix microhardness (about 188 HV0.1), the cubic structure can be presumed as ferrite [23]. Based on the globular ferritic microstructure this DSC sample shows parallels with the simulation (Fig. 1), which indicates an austenite into ferrite transformation at a temperature of nearly 800 °C for solidification and cooling rates near equilibrium. Furthermore, in microscopic images a small amount of carbides at the grain boundaries forming a partly connected network is recognizable (see inset of Fig. 2a). They are rich in V and were assumed as MC carbides with M = V and Cr (not shown here) but were not detectable in the XRD pattern. Due to the carbide morphology, their position at the grain boundaries as well as the results of previous studies, the carbides can be identified as MC type [18,19,26,27]. M7C3 carbides like displayed in the simulation (Fig. 1) were not found by applied techniques. In contrast, alloy specimens cooled down with 10–50 K/min (Fig. 2b–f) exhibit a dendritic structure. On the basis of the XRD pattern, exemplarily shown for the 20 K/min sample in Fig. 3b, the included crystal structures of the dendritic DSC specimens were identified. For the determination of their lattice parameters the present phases were indexed based on structure models according to Ref. [18]. The patterns predominant peaks represent a bcc crystal structure exhibiting a lattice parameter a = 0.28720(3) nm, which could refer to martensite. The significantly broadened XRD peaks compared to the ferritic specimen and slightly larger lattice parameter give hints for a lattice distortion and slight expansion due to the carbon dissolved in martensite [19]. Furthermore, the detected fcc structure can be assigned to an austenitic phase with a lattice parameter of a = 0.3601(4) nm. Beside these two main phases, there are signs in the XRD pattern indicating the existence of carbides. Due to their low content the carbides are at the detection limit, only their main peak is visible fairly poor, so they were neglected in the performed Rietveld analysis. As a result, the sample consists of 92 wt% martensite and 8 wt% austenite. Fig. 4 shows exemplarily the microstructure of the dendritic DSC sample cooled down with 20 K/min at higher magnifications. Hereby the dark brown structural constituents in the dendritic area are martensite plates. They are embedded in a bright appearing phase, which is composed of retained austenite. Furthermore a
small volume of carbides, visible as skeleton-like light gray phases, is situated at the former austenite grain boundaries. The carbides are clearly visible as dark phase in the SEM image in backscattered electrons (BSE)-contrast in Fig. 5a. Thereby two types of carbides with varying shapes and gray scales were observed indicating different chemical compositions. As their volume fraction was too low for a clear identification in the XRD pattern, the characterization of the carbides was done by EDX and EBSD. Based on the EDX results shown in Fig. 5b–e, two types of carbides can be distinguished: those that are rich in V and those that exhibit mainly Cr with notable amount of Fe. Typical EBSD diffraction patterns taken from the V-rich and Cr-rich carbides are displayed in Fig. 6. The experimental results are in accordance with the found indexed patterns showing the Miller indices of the zone axes. The V-rich carbides accord well with a fcc crystal structure referring to the MC-type. For the Cr-rich carbides a good accordance with the orthorhombic M7C3 structure was detected. The analyzed carbide types are in agreement with the results of the simulation. Apart from the carbides, single martensite plates are recognizable in the SEM image (Fig. 5a). In addition, the enrichment of the austenite with chromium can be seen in the EDX data (Fig. 5b–e) as reported in Ref. [19]. By means of the previous examinations the influence of the cooling rate on the evolving phases of FeCrVC is shown. Only slow cooling rates lead to a ferritic matrix according to the simulation displayed in Fig. 1, whereas at rates higher than 10 K/min a microstructure mainly composed of martensite and retained austenite is already formed. Because the cooling rate determines the evaluation of phases in the alloy, it is anticipated that the mechanical properties will likewise be influenced by cooling rate. This is confirmed by the results of the microhardness measurement in Table 3. A significant rise in the microhardness of the DSC samples between the samples cooled with 5 K/min from about 180 HV0.1 to more than 450 HV0.1 is observed. This is attributed to the change of the main phase from soft ferrite to harder martensite [23]. Thereafter DSC samples with proceeding increase of the cooling rate exhibit a continuously rising microhardness. It is assumed to be caused by an increasing strength of the martensitic phase. The reason for this could be increasing deformations in the martensite lattice due to higher contents of carbon and alloying elements [17]. This can be concluded from the simultaneous drop of the martensite transformation temperature, which is described below (see Section 3.3.). Beside the effect of the cooling rate on the occurring phases and resulting hardness, an impact on the morphology size is observed, too. The dendritic structure (Fig. 2) is continuously refined with increasing cooling rate, which was quantified by measuring the SDAS. Due to the small specimen size in relation to the size of the dendritic structure each DSC sample contained only few dendrites. Hence less than five SDAS could be captured whereupon the requirements for an accurate measurement according to Ref. [21] were fulfilled only partially. The result of the measured SDAS as arithmetic mean were plotted against the corresponding cooling rate during solidification in a double-logarithmic scale. Hereby SDAS k2 and cooling rate T_ were found to be related by the equation
k2 ¼ AT_ b : The parameters were determined as A = 285.9 lm and b = 0.398 for cooling rate values in K/min. The rising cooling rate is accompanied by a smaller solidification time. As the diffusion in the melt is higher than in the solid, this is followed by depression of the Ostwald ripening process eventually leading to smaller dendrites [28].
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Fig. 2. OM images of as-solidified FeCrVC-samples cooled with 3 (a), 10 (b), 20 (c), 30 (d), 40 (e) and 50 K/min (f), etched with Nital; inset in (a) shows SEM image of carbides along the grain boundary in FeCrVC-sample.
Fig. 3. Experimental XRD pattern of the DSC samples with cooling rates of 3 K/min (a) and 20 K/min (b) with identified phases, calculated fit (Rietveld method) and difference plot (difference of calculated and experimental curve).
3.3. Differential scanning calorimetry Prior to the cooling process the as-cast specimens were heated into the liquid state up to 1480 °C. The obtained DSC heating and cooling curves show several endothermic and exothermic peaks,
which are attributed to phase transformations in the specimen. The melting process is described by the endothermic peaks of the heating curve, whereas the cooling curve exhibits exothermic peaks. Fig. 7a shows a comparison of the heat flow signals for different cooling curves. The two DSC curves obtained for each
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Fig. 4. OM image of the as-solidified structure taken from the 20 K/min sample showing the dendritic morphology composed of martensite, austenite and carbides, etched with Nital.
cooling rate were found to be consistent, certifying the measured reaction sequences as reliable. Therefore, only one representative curve for each cooling rate was displayed in Fig. 7a. Due to the focus on the solidification behavior, a typical DSC cooling curve (cooling rate 20 K/min) is displayed in detail in Fig. 7b and c. The solidification part of the DSC curve in Fig. 7b shows two major peaks, whereas the first one is a double peak. The start of the solidification is characterized by the onset of the
Fig. 5. SEM image of the 20 K/min sample taken in BSE-contrast (a) with related EDX maps of iron (b), chromium (c), vanadium (d) and carbon (e).
first event, which represents the liquidus temperature Tl. It can be assumed that this first peak section can be attributed to the formation of primary delta ferrite. Although the simulation for the alloy composition (Fig. 1) does not predict a formation of delta ferrite, the typical delta ferrite peak [11] is visible for slow heating and cooling rates up to 20 K/min (Fig. 7a). This can be mainly attributed due to carbon escaping into the atmosphere during the DSC measurements. At slow heating and cooling rates, a larger quantity of carbon escapes, so that the phase formation calculated for an alloy with 1 wt% of carbon (Fig. 1) is not valid anymore and the formation of delta ferrite becomes possible. For faster DSC measurements (P30 K/min) less carbon can evaporate. This leads to a direct transformation of liquid into austenite and, therefore, no double peak is formed (as shown in Fig. 7b). To verify this assumption the final carbon content of several DSC samples was quantified by carrier gas hot extraction directly after the performed DSC measurements. For heating and cooling rates of 20 K/min the specimens contain in average 0.48 wt% carbon and for 40 K/min an average carbon content of 0.8 wt% was determined. These measurements show that the final carbon content is significantly influenced by the heating and cooling rate of the DSC sample and, thus, the amount of carbon escaping during the DSC measurements essentially affects the solidification path of the alloy. Fig. 8 shows the mass fraction of the equilibrium phases as a function of the temperature for a similar alloy where the carbon content is 0.5 wt% and the iron-content is slightly higher (Fe93.2Cr4.2V2.1C0.5) compared to the base alloy. It is shown that delta ferrite is formed, which is consistent with the results of the DSC measurements. It can be assumed that cooling rates slower than 20 K/min do have the same or even less carbon content. In comparison to microscopic results of the DSC sample cooled down with 3 K/ min, the calculations with 0.5 wt% C (Fig. 8) show more reliable formation pathes and phase fraction values for the arising carbides than the simulation for 1 wt% C (Fig. 1). After solidification and fully formation of an austenitic matrix carbides of the MC type form at the austenitic grain boundaries as shown in Fig. 2a. M7C3 carbides are generated relatively late in the solid state after the transformation of austenite into ferrite. The simulation with 0.5 wt% C (Fig. 8) depicts for the M7C3 carbides a phase fraction which is distinctly less than in the first calculation (Fig. 1). These carbides should be very fine and hardly detectable by XRD or microscopic measurements which was already assumed in Section 3.2. The second part of the double peak in Fig. 7b embodies the peritectic transformation of liquid and delta ferrite in austenite [11,25]. Hereafter, the second peak represents an eutectic reaction where the remaining melt is decomposed into austenite and MxC carbides [11,29], which have a MC structure as shown by microstructure analysis where the eutectic MC (V rich) carbides are embedded in austenitic phase along the grain boundaries (Fig. 4). In this study the offset of the second peak is taken as solidus temperature Ts. The further DSC cooling curve in the solid state is displayed in Fig. 7c. There are three exothermic peaks, which exhibit a much lower heat flow than the solidification peaks. It can be assumed that the peak at 1100 °C represents the precipitation of carbides with a M7C3 (Cr rich) structure found in the DSC sample. After the solidification was completed this carbide type could form out of the Cr enriched matrix on the energetically favoured interfaces of eutectic MC carbides in a very fine morphology as shown in Fig. 5a. A further minor peak arises around 700 °C, which is typical for the austenite-ferrite transformation. It is possible that a small amount of bainite structure has formed, which has to be checked by further examinations. The large peak between 300 and 400 °C can be attributed to the formation of martensitic phase, which is
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Fig. 6. Experimental EBSD patterns taken from the vanadium (a) and chromium rich carbides (c) fit with the indexed patterns of fcc-type VC (b) and orthorhombic Cr7C3 (d).
Table 3 Microhardness of the DSC sample cooled with 5 K/min measured in ferrite grains as well as measured microhardness in the dendritic area of the alloy mainly composed of martensite and retained austenite (10–50 K/min). Cooling rate (K/min)
5
10
20
30
40
50
Microhardness (HV0.1)
173 ± 40
456 ± 19
498 ± 13
594 ± 18
651 ± 28
748 ± 31
Fig. 7. DSC cooling curves of FeCrVC at various scanning rates in comparison (a) showing the position of liquidus Tl, solidus Ts and martensite start temperature Ms; DSC curve for a cooling rate of 20 K/min showing the solidification (b) and cooling in the solid state (c).
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Fig. 8. Mass fraction of equilibrium phases of Fe92.2Cr4.2V2.1C0.5 as a function of temperature, simulated with JMatPro.
contained in the 20 K/min DSC sample (Fig. 4). Hereby the peak onset is characterized as martensite start temperature Ms. In general the heat flow signal increases with higher cooling rates due to the rising enthalpy [8,12]. For this reason signals from faster cooling rates have a better visibility but also tend to overlap and conceal single reactions [30]. Thus the peaks for the delta ferritic and austenitic solidification can be clearly distinguished up to 20 K/min, while at higher rates only one major peak is visible. Apart from the solidification peak, signals with cooling rates of 3 and 5 K/min exhibit around 800 °C another evanescent peak, which is visible only in detail (not shown here). The DSC samples cooled with 3 and 5 K/min have a very low microhardness due to the formation of the ferritic phase. Hence the peak in the heat flow can be attributed to the transformation from austenite into ferrite, which was proposed by the simulation in the same range of temperatures. DSC curves with cooling rates of 10 K/min and above features another peak arising around 400 °C (Fig. 7a). This peak can be attributed to the formation of martensite, which is present in the mentioned alloy specimens. By examining the identical peaks at different cooling rates a general shift toward lower temperatures at higher cooling rates is observed. As a consequence, the characteristic temperatures Tl, Ts, and Ms varying with the cooling rate. Thereby the deviation of the martensite start temperature can be mainly assigned to the escaping carbon during the DSC measurements at slower heating and cooling rates. The results in Fig. 9 illustrates that an increasing cooling rate leads to lower liquidus and solidus temperatures. DSC heating curves do not show this dependency. Hereby the dependence between the liquidus and respectively solidus temperature and DSC cooling rate can be described as linear approximation. By linear extrapolation it was possible to estimate liquidus and solidus temperature at a zero velocity in equilibrium [13]. The values found are almost equal for the heating and cooling process as seen in Table 4. Furthermore, the experimentally determined temperatures were found to be in good agreement with values of the simulation and approximated calculations. Thereby the approximated liquidus temperature was determined by the empirical formula given in Ref. [31] via the chemical constitution of FeCrVC:
Fig. 9. Influence of the cooling and heating rate on the liquidus (a) and solidus temperature (b) of FeCrVC.
Table 4 Liquidus and solidus temperature of FeCrVC in equilibrium, linear extrapolation of data, simulation by JMatPro, approximation after formula in Ref. [31].
Tl (°C) Ts (°C)
DSC-heating
DSC-cooling
Simulation
Approximation
1456 1283
1453 1284
1445 1290
1455 –
However, the solidification time tf obviously decreases at higher cooling rates due to the accelerated cooling process. Hereby solidification time and cooling rate are related by an inverse proportionality. 4. Summary and outlook The present work demonstrates that the cooling rate has a distinctive influence on the microstructure and properties of the investigated alloy Fe92.7Cr4.2V2.1C1.
T l ð CÞ ¼ 1531—61:5C—4Mn—14Si—45S—30P—1:5Cr —2:5Al—3:5Ni—4V—5Mo ðcontents in wt%Þ: Altogether the observed drop of solidus temperature is larger than for the liquidus temperature. This leads to a higher solidification interval DT with increase of the cooling rate, which was found to rise linearly as presented in Fig. 10.
Fig. 10. Solidification time tf and interval DT of FeCrVC as a function of cooling rate.
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In this study, DSC samples were heated and solidified at various cooling rates between 3 and 50 K/min. With the help of the obtained alloy specimens and appropriate DSC signal curves, the influence of the cooling rate was systematically analyzed and characterized. The results show that the temperature control during DSC measurements significantly affects the solidification process. Liquidus and solidus temperature are lowered with increasing cooling rates whereas the solidification interval, which is important for the processing of the material, is expanding. The estimated equilibrium liquidus and solidus temperatures are 1453 °C and 1284 °C respectively. Furthermore, an influence on the evolving phases and morphology was observed. Specimens manufactured at slow cooling rates solidify in a cellular structure and consist mainly of ferrite. In contrast, specimens cooled at higher rates exhibiting dendritic microstructure composed of the non-equilibrium phases martensite and retained austenite. Simultaneously a refinement of the dendritic morphology as well as a depression of the martensite start temperature with an increasing cooling rate was observed. Beside the main phases, a network of mixed carbides was detected and verified as MC (M = V, Cr) and M7C3 (M = Cr, Fe). Moreover, the investigated alloy exhibits a rising microhardness at increasing cooling rates. Altogether, the presented work clarifies the connection between cooling rate, secondary dendrites arm spacing and resulting microhardness for a newly developed tool steel. This allows to draw first conclusions from the solidification parameters with regard to the resulting mechanical properties. Thereby it could be displayed that the carbon content is significantly influenced by the temperature regime and therefore affects the phase transformations. The gained knowledge is a step toward systematically adjusting and enhancing the properties of as-solidified tool steels by tailored temperature control. The equilibrium phase simulation consulted for comparison provided good results regarding the solidification parameters and occurring phases. In future the simulation results will be combined with the displayed parameter relations. This should provide the opportunity for more reliable simulations in the non-equilibrium state and offers novel possibilities for predictions close to real production processes. Acknowledgement The authors would like to thank B. Bartusch, H. Bußkamp, S. Donath, M. Frey, C. Geringswald, W. Gruner, V. Hammond, R. Keller,
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