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Failure modes in field-tested brass die casting dies
Journal of Materials Processing Technology 148 (2004) 108–118
Failure modes in field-tested brass die casting dies Anders Persson a,∗ , Sture Hogmark b , Jens Bergström a a
Department of Materials Engineering, Karlstad University, SE-651 88 Karlstad, Sweden b The Ångström Laboratory, Uppsala University, SE-751 21 Uppsala, Sweden
Received 30 June 2003; received in revised form 30 June 2003; accepted 15 January 2004
Abstract Tools for die casting of, for example, brass and aluminium alloys are exposed to severe thermal, mechanical, and chemical conditions. The performance and service life of the die components are limited because of reasons such as thermal fatigue cracking, erosion, corrosion, soldering, and gross fracture. To minimise these damages, the dies are normally made of hot work tool steel. This study aims at elucidating the life-limiting failure mechanisms in dies aimed for brass die casting. Two cavity inserts and eight cores of two hot work tool steels, quenched and tempered to different conditions, were examined and evaluated with respect to failure mechanisms after use in actual brass die casting. It was found that the dominating failure mechanism in the investigated tools was thermal fatigue cracking. The thermal fatigue crack initiation is associated to accumulation of the local plastic strain that occurs during each casting cycle, typical of a low-cycle fatigue process. The initial growth of the thermal cracks is facilitated by oxidation of the crack surfaces, and proceeded growth is facilitated by additional oxidation in combination with crack filling of cast material, and by softening of the tool material. The most striking observation was a mechanism of crack growth promotion that involves a local enrichment of Pb from the brass alloy melt at the crack front. Consequently, in addition to increasing the overall tool steel yield strength and/or oxidation resistance, there is a potential to improve the tool life either by removing or substituting the Pb in the brass casting alloy with any non-harmful element, or using a tool material not susceptible to Pb embrittlement. © 2004 Published by Elsevier B.V. Keywords: Failure mechanisms; Thermal fatigue; Hot work tool steel; Die casting; Field test; Brass
1. Introduction Die casting of brass is an industrially important method of forming near net-shaped components with high mechanical properties and corrosion resistance for heating, ventilation, and sanitary installations, for example valves, pipe couplings, etc. [1–3]. Prior to casting, the die is normally preheated to a temperature within the range of 300–350 ◦ C. During a die casting cycle, molten metal is forced into an internally cooled mould by the application of pressure. The peak metal pressure during the injection can exceed 70 MPa. A distinguishing characteristic of the process is that the liquid metal flows with high velocity during injection and provides rapid filling of the die cavity, typically within milliseconds. The entrance velocity of the melt is of the order of 1–10 m/s. The high velocity is necessary to fill the mould of thin-walled and complex-shaped products. Continuous internal cooling of the die during ∗ Corresponding author. Tel.: +46-54-700-1821; fax: +46-54-700-1449. E-mail address: [email protected] (A. Persson).
0924-0136/$ – see front matter © 2004 Published by Elsevier B.V. doi:10.1016/j.jmatprotec.2004.01.052
the process makes the solidification of the casting efficient and, as a consequence, high rate manufacturing of typically 100 castings/h is enabled. When the casting has solidified, the die is opened and the casting is ejected. Subsequently, the die may be externally cooled and lubricated by spraying. A die casting mould is composed of several tool components, for example main die, cavity inserts, cores, and ejector pins [1–3]. Wear and failure of die casting dies has a technical and economical significance, since it degrades the surface finish and changes the dimensions of the tool and, therefore, also those of the casting, and it may even cause expensive failures of the tool. To withstand the damages during service the dies for brass die casting are frequently made of hot work tool steels, such as AISI H13, H20, H21, or H22. They are used in quenched and tempered conditions with a hardness range of 360–520 HV, depending on die component and steel grade. High levels of hot yield strength, temper resistance, ductility, and toughness are important mechanical properties of these tool materials.
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1.1. Failure modes in die casting dies
2. Experimental
Thermal fatigue cracking is the most important life-limiting failure mode in dies for brass die casting, since the temperature of the melt is high (about 970 ◦ C) [1–4]. The cracking results from rapid fluctuations in temperature, stress, and strain in the die surface because of the cyclic nature of the casting process. The temperature cycling may induce stresses high enough to impose an increment of plastic strain in the tool surface during each cycle. Surface cracks appear already after a few thousand castings, or even earlier, and are, therefore, formed in the low-cycle fatigue range [5,6]. It is reported that creep and oxidation may significantly contribute to the cracking [5–9]. The thermal fatigue cracking is often observed on the tool surface as a network of fine cracks, often named heat checking. Another variant of thermal fatigue cracks (stress cracks) may be observed as individual and clearly pronounced cracks in areas exposed to local stress concentrations [5]. The formation of thermal fatigue cracks may lead to loss of surface material in the form of small fragments (splintering). In addition to thermal fatigue cracking, erosion, corrosion, local adherence of the casting alloy to the tool (soldering), and gross fracture are other important failure modes. Erosion is induced by the high velocity of the incoming melt and is partially aggravated by the presence of solid particles in the molten casting alloy. The erosive damage is primarily seen where the molten metal jet first hits the die surface. Corrosion damage originates from dissolution of the tool material into the liquid metal. Erosion and corrosion may cause significant loss of surface material, and may also increase the tendency of soldering. Soldering results from the interaction between the die material and the casting alloy during injection and solidification and may be observed as residuals of the casting alloy on the die surface after ejection. Gross fracture is primarily caused by excessive thermal shocks or mechanical overloading, and leads often to total failure of the tool. All modes of tool failure limit the performance and life of die casting tools, which, ultimately, ends in higher production costs and/or environmental impacts. Only a few studies on failure mechanisms in field-tested die components used in brass die casting can be found in the literature [6,7], while those on aluminium die casting are more frequent [10–12]. The significantly higher melt temperature in die casting of brass, the lower entrance velocity of the melt, the chemical properties of the melt as well as other process parameters clearly separate the casting conditions of brass and aluminium. This motivates the present study, which aims at elucidating the life-limiting failure mechanisms in field-tested brass die casting dies. Two cavity inserts and eight cores were examined. They were made of two hot work tool steels, quenched and tempered to different conditions, and evaluated with respect to hardness changes, growth and density of cracks, and other failure mechanisms.
2.1. Materials
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The two cavity inserts were made of a hot work tool steel, QRO 90 Supreme (Uddeholm Tooling Designation), with the nominal chemical composition (wt.%) of 0.38 C, 0.39 Si, 0.75 Mn, 2.6 Cr, 2.25 Mo, 0.9 V, and Fe balance. The eight cores were made of two different hot work tool steels, QRO 90 Supreme and Hotvar (Uddeholm Tooling Designation). The nominal chemical composition (wt.%) of Hotvar is 0.55 C, 1.0 Si, 0.75 Mn, 2.6 Cr, 2.25 Mo, 0.85 V, and Fe balance. The cavity inserts and cores were hardened and tempered, according to Table 1. The brass alloy, CuZn33Pb2Si–C (Ametal C, Tour and Andersson Designation), with an approximate chemical composition (wt.%), as obtained on a casting by X-ray spectroscopy of 64.1 Cu, 32.3 Zn, 1.9 Pb, 0.71 Si, 0.35 Fe, 0.05 As, 0.03 Al, was used as casting material throughout the field tests. The alloy has a nominal liquidus temperature of 887 ◦ C and a solidus temperature of 844 ◦ C. 2.2. Die casting conditions for the investigated tools All cavity inserts and cores have been used for production of brass valves. The geometry of the studied die components is shown in Fig. 1. The field tests were made in a 2.8 MN cold chamber machine. The temperature of the brass melt was 970 ◦ C and the total cycle time for one casting was 48 s. The entrance velocity of the melt was 1–2 m/s and the maximum metal pressure during injection was approximately 190 MPa. The total shot weight of each casting was 2.1 kg. Water at 20–50 ◦ C was continuously circulated through cooling channels in the cavity insert. After ejection of each casting, the surfaces of the tools were lubricated but not intentionally cooled by spraying a commercial lubricant fluid (Klüber Metalstar Fe 143). The cavity inserts and cores were shot-blasted after each 600 and 200 castings, respectively, and also after the tests. This treatment was performed to remove deposits of Zn compounds on the tool surface and to decrease the tendency of soldering. Note that the cavity inserts in this study were not preheated. The two studied cavity inserts and two of the cores had been used in production until the surface finish or the dimensional changes, or both, of the castings exceeded the tolerances. The tool lifetime of the studied cavity inserts Table 1 Nominal heat treatment and hardness of the field-tested cavity inserts and cores Die component
Steel grade
Austenitising (◦ C/min)
Tempering (◦ C/h)
Cavity insert Cavity insert Core Core
QRO 90 QRO 90 QRO 90 Hotvar
1020/35 1020/34 1020/57 1050/17
640/2 615/2 625/2 580/2
× × × ×
2 2 2 2
Hardness (HV) 430–440 480–500 460–480 600–620
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Fig. 1. The field-tested tool components used in brass die casting. (a) Cavity insert: the arrow indicates the gate where the melt enters. (b) The studied core, located closest to the gate. A and B indicate the positions of the detailed evaluation of the cavity inserts and cores, respectively. The diameter of the cavity at A is 33 mm and that of the core at B is 29 mm.
was 45 000–50 000 castings, whereas that of the cores was 10 000–15 000 castings. Of the eight studied cores, one of each material was investigated after their expired lifetime, whereas one core of each material was taken out of production for examination after 500, 1000, and 1500 castings, respectively. 2.3. Characterisation Identification of the failure mechanisms was performed by macroscopic examinations as well as by detailed studies on the surface of the tools, and on fractured and polished cross-sections, using light optical microscopy (LOM), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS). The fractured cross-sections were obtained after having cooled the steel in liquid nitrogen. For the cavity inserts and cores, the detailed evaluations was focused to an area which was expected to be representative of thermal fatigue cracking with a minimum of stress concentrations (see Fig. 1). The core located closest to the gate was examined since the casting conditions are expected to be most severe where the melt enters the cav-
ity (see Fig. 1b). For example, the melt temperature is at its highest level, the time of contact between the casting alloy and the tool is the longest, and the erosive and corrosive damages resulting from the incoming melt is expected to be most distinguished on that core. Hardness versus depth profiles to reveal any thermal softening were obtained by Vickers indentations on polished cross-sections, using a load of 25 g. The evaluation of the cavity inserts and cores, respectively, were performed at the same location in each die component to minimise any dependence on differences in the casting conditions in the tool.
3. Results 3.1. Macroscopic failure mechanisms in field-tested brass die casting dies The macroscopic surface damages on the two worn-out cavity inserts and cores, respectively, were severe and observed as a network of surface cracks and as individual and clearly pronounced cracks at sharp corners (see Fig. 2). Note
Fig. 2. Typical macroscopic surface damages observed on worn-out cavity inserts and cores. (a) Overview of an insert revealing local severe surface cracking (A) and individual cracks at sharp geometrical corners (B). (b) Overview of a core revealing a severe network of surface cracks, individual cracks at sharp radiuses, and remnants of solidified brass within the cracks (indicated by the arrow).
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Fig. 3. Macroscopic damages observed on cores. (a) Overview after 500 castings revealing individual and clearly pronounced cracks at sharp edge (indicated by the arrow). (b) Overview of a worn-out core revealing a network of surface cracks, and a gross fracture (indicated by the arrow).
the locally severe surface cracking on the worn-out insert, where a protuberance in the tool has been contacted by the melt from several directions. On one of the worn-out cavity inserts, it was observed that some cracks had grown together and caused loss of small steel fragments. In general, the crack networks on the worn-out inserts are not as rough as those on the worn-out cores. Macroscopic inspections revealed local filling of solidified brass in the cracks on the worn-out cores. Typically, individual and clearly developed cracks as well as local filling of solidified brass in the cracks were detected on the cores already after 500 castings (see Fig. 3a). The two worn-out cores showed these defects in a further pronounced stage, in addition to a relatively coarse crack network (see Fig. 3b). Gross fracture was observed on one of the worn-out cores, evident by the missing tip in Fig. 3b. The network of cracks on the two worn-out cores consisted of both wide and narrow cracks, locally filled with solidified brass (see Fig. 4). Note that wide cracks form the network, while more narrow cracks seem to grow into or within the areas separated by the wide cracks (also seen in Fig. 2b). The tendency of brass filling the cracks seems to
Fig. 4. Typical surface damages observed on worn-out cores (SEM). A severe network of wide and narrow cracks and remnants of brass within the cracks are revealed. The vertical scratches on both the surface of the tool (A) and on the solidified brass (B) were made after tool disassemblement.
be larger for the wide than for the narrow cracks. No brass is adhered to the areas separated by the cracks. 3.2. Thermal fatigue cracks 3.2.1. Fractured cross-sections It proved possible to reveal the geometrical details of the cracks by making fractured cross-sections, since the surface of a thermal fatigue crack was easily distinguished from the imposed fracture. At an early stage, the cracks are semi-circular (see Fig. 5). Later on, the crack may grow to a considerable depth, see Fig. 6a, and also extend in the lateral direction (see Fig. 7). The position of the crack front prior to the imposed fracture coincides with a strong concentration of Pb as detected by EDS. Lead also shows up brighter in the micrographs due to atomic number contrast (see Fig. 6a and b). At the crack front of Fig. 6a, a small semi-elliptic sub-crack was detected (see Fig. 6b). A dimple pattern is revealed along the crack tip with bright contrast (see Fig. 6c). 3.2.2. Polished cross-sections Studies of polished cross-sections further elucidate the morphology of the crack pattern (see Fig. 8). The cracks tend to grow perpendicular to the tool surface and reach several
Fig. 5. Fractured cross-section of a core after 500 castings revealing a large semi-circular crack (SEM). The original fatigue crack front is indicated by the arrow.
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Fig. 6. Fractured cross-sections of worn-out cores (SEM). (a) Overview revealing a major crack and a small sub-crack initiated at the front of the major crack (indicated by the arrow). The frame indicate EDS-mapped area, see below. (b) Close-up of the front of the sub-crack of (a). (c) A further increase in magnification reveals, through the characteristic dimples, the presence of a ductile material at the crack tip. (d) EDS map of Pb of the sub-crack of (a). High concentration is represented by bright areas.
millimetres in depth (see Fig. 8a). All of them are filled with a material that has a layered structure next to the steel walls, and a more inhomogeneous structure in the centre where it contains voids and cracks (see Fig. 8b). The appearance of a typical crack tip and the material filling it is seen in Fig. 9. The only element detected at the very crack tip was Pb (see Fig. 9b). Lead also reveals itself in Fig. 9a due to its high atomic number contrast. Most of
Fig. 7. Fractured cross-section of worn-out core revealing a major fatigue crack with the crack front (indicated by the arrow) extending parallel to the surface of the tool (SEM).
the material filling the crack also contains O, indicating a strong element of oxidation (cf. Fig. 9d). The cracks are typically filled already at an early stage of development, and especially after further development (see Fig. 10). Figs. 11 and 12 reveal that their surfaces mainly consist of iron oxides and that the interior is filled with residuals from solidified brass. Pb was not detected in the tip area of this crack. Relatively short cracks (cf. Fig. 13), filled with a mixture of iron and zinc oxides, and a diffuse Cr- and O-rich layer in between the oxides filling the crack and the tool steel, were also frequently observed. Neither Cu nor Pb was detected in these cracks. In general, the appearance of polished cross-sections of the cavity inserts resembles those of the cores. 3.2.3. Crack length and density The maximum length, mean length, and density of the thermal fatigue cracks extracted from polished cross-sections of the tools are of the order of 0.4–3.2 mm, 0.03–1.5 mm, and 0.5–5.5 mm−1 , respectively, for all die components at all number of castings produced (see Fig. 14). After the initial 500 castings, the maximum and mean crack length in the cores is shorter in the harder than in the softer tool, whereas the crack density is higher in the harder tool
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Fig. 8. Polished cross-section of a worn-out core (SEM). (a) Overview revealing typical thermal fatigue cracks filled with material. (b) Close-up of the crack indicated by the arrow in (a).
steel (see Fig. 14). Evidently, the difference in maximum and mean crack length and crack density in the cores averages with number of castings, and reach approximately the same level when they are worn-out, independent of steel grade and initial condition. For the worn-out cavity inserts, the maximum crack length is somewhat shorter and the crack density is significantly lower in the harder than in the softer tool, whereas the mean crack length shows the inverse correlation with hardness. Finally, it is clearly seen that the maximum as well as the mean crack length is significantly shorter in the worn-out cavity inserts than those in the worn-out cores. Evidently, the crack density in
the worn-out die components is higher in the cavity inserts than that in the cores. 3.3. Hardness profiles As expected, for all die components the surface hardness is decreasing with the number of castings produced (see Fig. 15). After the first 500 castings, the hardness reduction of the cores is limited to a surface layer of about 2 mm, and the loss of hardness is larger for the harder than for the softer tool (see Fig. 15a and b). Evidently, the difference in surface hardness between the cores is reduced and the
Fig. 9. Polished cross-section through a typical crack tip in high magnification. (a) Compositional SEM image illustrating the layered filling. (b–d) EDS maps. The element concentration is proportional to the brightness.
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(cf. Fig. 6a, b, and d). Obviously, Pb from the brass casting alloy concentrates at the crack front during casting and due to interaction with the tool steel it is able to locally assist the crack growth. Since the thermal cracking failure dominates, further consideration on this mechanism is given below. 4.1. Thermal cracking
Fig. 10. Polished cross-section of a worn-out core revealing typical long and short cracks (LOM). The frames indicate EDS-mapped areas, see below.
softened surface layer increases in thickness with number of castings. When the cores are worn-out, the softened surface layer extends beyond 10 mm depth. For the worn-out cavity inserts, it is observed that the hardness in the superficial surface layer is reduced to approximately the same level, whereas an approximately constant difference in hardness between the two tools is maintained deeper below the surface (see Fig. 15c). The softened surface layer of the two cavity inserts is limited to a depth of less than 1.5 mm. 4. Discussion It was observed that the dominating failure mechanism in the investigated brass die casting cavity inserts and cores was thermal fatigue cracking (cf. Fig. 2). One of the most striking results from this investigation was that the thermal fatigue crack front advancement is associated with a local enrichment in Pb, whereas it is pinned in areas without Pb
4.1.1. Crack nucleation and growth Thermal cracks are nucleated when the local tensile stress in the surface exceeds a critical level. The tensile stresses are, in turn, a result of plastic yield of the tool material during the hot phase of the casting cycles. It was found that it takes less than 500 brass die casting cycles to nucleate and propagate the first thermal fatigue cracks in the cores (cf. Figs. 3a and 14). Consequently, the temperature alternations in the tool surface layer are severe enough to impose compressive stresses high enough to cause local accumulation of plastic strains to nucleate and propagate thermal cracks in the tool surface by a low-cycle fatigue process. The initial growth of the thermal fatigue cracks is facilitated by an oxidation attack on the crack surfaces to form a Cr- and O-rich layer in between oxides filling the crack and the visually unaffected tool steel (cf. Fig. 13 and Refs. [6,7]). Presence of oxides and brass materials in the cracks increases the compressive stresses and plastic yield during the hot phase of the thermal cycle. The result is a further increase in tensile stress during the cold phase, which results in crack growth. The thermal cracks advance considerably in length while they are typically filled with a mixture of oxides [6], and residuals from the brass alloy [7] (cf. Figs. 10–12). The oxides filling the cracks consist mainly of iron oxides [7], but also of zinc oxides and lead oxides [6], as shown in Figs. 9–13. Their structure in between the steel walls are layered or inhomogeneous and contain cracks
Fig. 11. EDS maps of the area indicated by frame A in Fig. 10. The element concentration is proportional to the brightness.
Fig. 12. EDS maps of the area indicated by frame B in Fig. 10. The element concentration is proportional to the brightness.
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Fig. 13. Polished cross-section through a typical crack of a worn-out core in high magnification: (a) LOM image; (b–f) EDS maps. The element concentration is proportional to the brightness.
and voids [6] (cf. Figs. 8 and 9). The layered structure of the filling material indicates that the crack filling aggravates gradually during the casting process. The tensile stresses imposed on the tool material surface layer during the cooling phase effectively cause local rupture of filled material and open the thermal cracks (cf. Figs. 8b and 13). These cracks in the filling material act as channels for more brass materials to fill the cracks and for oxygen to penetrate down to the crack tip area and oxidise the steel and also the filling material (cf. Figs. 8–12). In addition, the opening of the cracks in the filling material gives an indication of the maximum amount of brass that can fill the crack during the next casting cycle (cf. Fig. 8b). 4.1.2. Crack growth promoted by Pb embrittlement Steel exposed to liquid lead may form solid Fe–liquid Pb systems that causes liquid metal embrittlement [13]. Embrittlement may also occur when the temperature is below the melting point of the embrittling metal (solid metal-induced embrittlement), but the severity of the embrittlement increases considerably at the melting point of the embrittling metal. The mechanism of liquid metal embrittlement is probably associated with a reduction of the atomic bond strength of the solid metal. Previously, temperature recordings in a 5 mm surface layer of a cavity insert during actual die casting of brass revealed that the maximum tool temperature was about 828–450 ◦ C, decreasing from the tool surface and inwards, and that the minimum temperature through the surface layer was about 300 ◦ C [14]. Thus, the temperature on the tool surface layer exceeds the melting point of Pb (327 ◦ C) during each casting cycle and results, consequently, in al-
ternating melting of Pb in the brass alloy filling the cracks (cf. Fig. 10). In this study, it was observed that in a relatively early stage, the thermal fatigue cracks have generated a semi-circular shape (see Fig. 5). Subsequently, the crack front expands in the areas where the imposed stresses are high enough to enable propagation (cf. Fig. 6a). The alternating melting of Pb during each casting cycle cause the liquid Pb from the brass alloy to penetrate through the crack down to the very crack tip, where it is enriched (cf. Figs. 6 and 9b). The enrichment in Pb at the very crack tip is supported by the absence of O in this area (cf. Fig. 9b and d), and also by the observation of a ductile material (Pb) at the crack front (see Fig. 6c). It is obvious that the crack front advances locally where there is an enrichment in Pb, whereas it is pinned in areas without Pb (see Fig. 6a, b, and d). Obviously, the chemical interaction with Pb from the brass alloy and the tool steel is able to locally assist the crack growth, most probably by liquid metal embrittlement. 4.1.3. Crack growth promoted by thermal softening of the tool steel The crack evaluations after the initial 500 castings revealed that the growth of thermal cracks are suppressed by higher tool hardness levels, as evidenced by the shorter maximum and mean crack lengths in the initially harder Hotvar core than in the softer QRO 90 core (see Fig. 14a and b). Thereafter, any differences in the crack characteristics among the cores averages with the number of castings to approximately equal levels at their expired tool life of 10 000–15 000 castings, independent of tool steel and initial condition.
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Fig. 14. Maximum crack length (a), mean crack length (b), and crack density (c) vs. number of castings in the cavity inserts and cores, based on cracks larger than ∼5 m. Each pile is based on measurements along a line of about 17 mm. The designations I and II correspond to cores of QRO 90 (460–480 HV) and Hotvar (600–620 HV), respectively, III and IV to cavity inserts of QRO 90 at nominal hardness values prior to the field tests of 430–440 and 480–500 HV, respectively.
The fact that the average of the crack characteristics with the number of castings (Fig. 14a and b) correlate strongly to the average of the surface layer hardness among them (Fig. 15a and b) indicates that the crack growth is facilitated by the gradual softening of the tool steels. 4.1.4. Crack network development and final appearance Already after 500 castings, individual and clearly developed cracks and crack filling were clearly detected on the cores by the naked eye (cf. Fig. 3a). Subsequently, a network of surface cracks and more individual cracks are formed and the tendency of crack filling aggravates (cf. Figs. 3b and 4). Naturally, the tendency of crack filling is low when the cracks are narrow and increases as they become wider (see Fig. 4). Eventually, some surface cracks may grow together and cause local detachment of surface material (splintering), or deeply penetrate into the tool and lead to gross fracture (cf. Fig. 3b). The observation that the worst surface cracking appeared in tool areas that are exposed to the most critical thermal
Fig. 15. Hardness vs. depth profiles after different number of castings: (a) cores of QRO 90 (460–480 HV); (b) cores of Hotvar (600–620 HV); (c) worn-out cavity inserts of QRO 90.
and/or mechanical stress concentrations, such as sharp corners and protuberances, or parts of the tool that had been more or less surrounded by the melt (cf. Fig. 2), is in agreement with previous findings [5,7]. The fact that the cores are exposed to more intense thermal conditions than the inserts is supported by their more pronounced hardness reduction demonstrated in Fig. 15. Further support is given by the observation that the thermal cracks penetrate considerably deeper into the cores than into the inserts (see Fig. 14a and b). 4.2. Influence of tool maintenance on crack characteristics Previously, it has been reported that the development of thermal fatigue cracks can be characterised by the three consecutive stages: crack nucleation, rapid crack growth, and retarded crack development [5,15,16]. After some initial cycles during which most cracks are formed, some cracks continue to grow deeper into the tool material [5]. Since areas along each thermal crack are locally stress relieved, the growth of adjacent cracks is retarded as the larger once propagate into the tool material [17].
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The tool maintenance applied during the die casting service conditions includes lubrication after each casting and intermittent cleaning by shot-blasting as well as other mechanical procedures, as indicated by the scratches on the tool surface in Fig. 4. From Fig. 14, it is obvious that the crack development in the cores neither can be perfectly characterised by the three consecutive stages, nor by the crack density saturation after the initial castings during which almost all cracks are formed. Surface material removal during tool maintenance is most likely the reason why the crack density in the cores has a tendency to decrease within the whole range of castings produced (see Fig. 14c). Another effect from this maintenance is that the crack length values of Figs. 14a and b are somewhat shorter than they certainly would have been without maintenance. The maximum crack length in the inserts (about 1.3–1.4 mm) after 45 000–50 000 castings (Fig. 14a) is in good agreement with the maximum crack length (about 1.5 mm) estimated from temperature recordings in a cavity insert during actual die casting of brass [14]. The cavity insert, tool steel, brass alloy, and the die casting process parameters during the temperature recordings resemble those of this study. It is therefore concluded that the maintenance has not reduced these maximum crack length values considerably. Finally, the tool maintenance is probably the reason why no macroscopic evidence of erosive or corrosive damage (including soldering) was identified on the investigated tools. However, those modes of tool failure appear to be of secondary importance in the degradation of the investigated tools. It is important to realise that if the field tests had been performed without any tool maintenance, it had not taken many castings before practical casting operation problems arise.
5. Conclusions In this study, the life-limiting failure mechanisms in field-tested cavity inserts and cores for die casting of brass were examined. The following conclusions can be drawn: • The dominating failure mechanism in brass die casting tools is thermal fatigue cracking. • The crack nucleation is associated to accumulation of the local plastic strain in the surface, typical of a low-cycle fatigue process. • The initial growth of the thermal fatigue cracks is facilitated by oxidation of the crack surfaces. • Proceeded growth is facilitated by filling of cracks with brass, by additional oxidation, and by softening of the tool material. • The oxides consist mainly of iron oxides, but also of zinc and lead oxides.
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• The most striking observation is a local enrichment of Pb at the crack front. • Since the crack has advanced locally in these regions, this Pb enrichment is suggested as an important mechanism of crack growth promotion in brass die casting. • In addition to increasing the overall tool steel yield strength and/or oxidation resistance, there is a potential to improve the tool life either by removing or substituting the Pb in the brass casting alloy with any non-harmful element, or using a tool material not susceptible to Pb embrittlement.
Acknowledgements The authors are grateful to Tour and Andersson AB and Mr. Jan Andersson at this company, and to Uddeholm Tooling AB and Bodycote Heat Treatment AB. The financial support from the Swedish Knowledge Foundation is also acknowledged. Special thanks to Dr. Richard Westergård at the Ångström Laboratory for performing the EDS mapping.
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Exposition, North American Die Casting Association, Cincinnati, 2001, p. 379. [13] B. Craig, ASM Handbook, vol. 13, 9th ed., ASM International, Metals Park, OH, 1987, p. 145. [14] A. Persson, On tool Failure in Die Casting, Doctoral dissertation, Uppsala University, 2003. [15] P.W. Hockanadel, G.R. Edwards, M.C. Maguire, M.D. Baldwin, The effect of microstructure on the thermal fatigue resistance of investment cast and wrought AISI H-13 hot work die steel, in:
Transactions of the 18th International Die Casting Congress and Exposition, Indianapolis, 1995, p. 343. [16] H.P. Lieurade, A. Dias, J. Giusti, A. Wisniewski, Experimental simulation and theoretical modelling of crack initiation and propagation due to thermal cycling, High Temp. Technol. 8 (1990) 137. [17] H.G. Baron, B.S. Bloomfield, Resistance to thermal stress fatigue of some steels, heat-resisting alloys, and cast irons, J. Iron Steel Inst. 197 (1961) 223.
Failure modes in field-tested brass die casting dies Anders Persson a,∗ , Sture Hogmark b , Jens Bergström a a
Department of Materials Engineering, Karlstad University, SE-651 88 Karlstad, Sweden b The Ångström Laboratory, Uppsala University, SE-751 21 Uppsala, Sweden
Received 30 June 2003; received in revised form 30 June 2003; accepted 15 January 2004
Abstract Tools for die casting of, for example, brass and aluminium alloys are exposed to severe thermal, mechanical, and chemical conditions. The performance and service life of the die components are limited because of reasons such as thermal fatigue cracking, erosion, corrosion, soldering, and gross fracture. To minimise these damages, the dies are normally made of hot work tool steel. This study aims at elucidating the life-limiting failure mechanisms in dies aimed for brass die casting. Two cavity inserts and eight cores of two hot work tool steels, quenched and tempered to different conditions, were examined and evaluated with respect to failure mechanisms after use in actual brass die casting. It was found that the dominating failure mechanism in the investigated tools was thermal fatigue cracking. The thermal fatigue crack initiation is associated to accumulation of the local plastic strain that occurs during each casting cycle, typical of a low-cycle fatigue process. The initial growth of the thermal cracks is facilitated by oxidation of the crack surfaces, and proceeded growth is facilitated by additional oxidation in combination with crack filling of cast material, and by softening of the tool material. The most striking observation was a mechanism of crack growth promotion that involves a local enrichment of Pb from the brass alloy melt at the crack front. Consequently, in addition to increasing the overall tool steel yield strength and/or oxidation resistance, there is a potential to improve the tool life either by removing or substituting the Pb in the brass casting alloy with any non-harmful element, or using a tool material not susceptible to Pb embrittlement. © 2004 Published by Elsevier B.V. Keywords: Failure mechanisms; Thermal fatigue; Hot work tool steel; Die casting; Field test; Brass
1. Introduction Die casting of brass is an industrially important method of forming near net-shaped components with high mechanical properties and corrosion resistance for heating, ventilation, and sanitary installations, for example valves, pipe couplings, etc. [1–3]. Prior to casting, the die is normally preheated to a temperature within the range of 300–350 ◦ C. During a die casting cycle, molten metal is forced into an internally cooled mould by the application of pressure. The peak metal pressure during the injection can exceed 70 MPa. A distinguishing characteristic of the process is that the liquid metal flows with high velocity during injection and provides rapid filling of the die cavity, typically within milliseconds. The entrance velocity of the melt is of the order of 1–10 m/s. The high velocity is necessary to fill the mould of thin-walled and complex-shaped products. Continuous internal cooling of the die during ∗ Corresponding author. Tel.: +46-54-700-1821; fax: +46-54-700-1449. E-mail address: [email protected] (A. Persson).
0924-0136/$ – see front matter © 2004 Published by Elsevier B.V. doi:10.1016/j.jmatprotec.2004.01.052
the process makes the solidification of the casting efficient and, as a consequence, high rate manufacturing of typically 100 castings/h is enabled. When the casting has solidified, the die is opened and the casting is ejected. Subsequently, the die may be externally cooled and lubricated by spraying. A die casting mould is composed of several tool components, for example main die, cavity inserts, cores, and ejector pins [1–3]. Wear and failure of die casting dies has a technical and economical significance, since it degrades the surface finish and changes the dimensions of the tool and, therefore, also those of the casting, and it may even cause expensive failures of the tool. To withstand the damages during service the dies for brass die casting are frequently made of hot work tool steels, such as AISI H13, H20, H21, or H22. They are used in quenched and tempered conditions with a hardness range of 360–520 HV, depending on die component and steel grade. High levels of hot yield strength, temper resistance, ductility, and toughness are important mechanical properties of these tool materials.
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1.1. Failure modes in die casting dies
2. Experimental
Thermal fatigue cracking is the most important life-limiting failure mode in dies for brass die casting, since the temperature of the melt is high (about 970 ◦ C) [1–4]. The cracking results from rapid fluctuations in temperature, stress, and strain in the die surface because of the cyclic nature of the casting process. The temperature cycling may induce stresses high enough to impose an increment of plastic strain in the tool surface during each cycle. Surface cracks appear already after a few thousand castings, or even earlier, and are, therefore, formed in the low-cycle fatigue range [5,6]. It is reported that creep and oxidation may significantly contribute to the cracking [5–9]. The thermal fatigue cracking is often observed on the tool surface as a network of fine cracks, often named heat checking. Another variant of thermal fatigue cracks (stress cracks) may be observed as individual and clearly pronounced cracks in areas exposed to local stress concentrations [5]. The formation of thermal fatigue cracks may lead to loss of surface material in the form of small fragments (splintering). In addition to thermal fatigue cracking, erosion, corrosion, local adherence of the casting alloy to the tool (soldering), and gross fracture are other important failure modes. Erosion is induced by the high velocity of the incoming melt and is partially aggravated by the presence of solid particles in the molten casting alloy. The erosive damage is primarily seen where the molten metal jet first hits the die surface. Corrosion damage originates from dissolution of the tool material into the liquid metal. Erosion and corrosion may cause significant loss of surface material, and may also increase the tendency of soldering. Soldering results from the interaction between the die material and the casting alloy during injection and solidification and may be observed as residuals of the casting alloy on the die surface after ejection. Gross fracture is primarily caused by excessive thermal shocks or mechanical overloading, and leads often to total failure of the tool. All modes of tool failure limit the performance and life of die casting tools, which, ultimately, ends in higher production costs and/or environmental impacts. Only a few studies on failure mechanisms in field-tested die components used in brass die casting can be found in the literature [6,7], while those on aluminium die casting are more frequent [10–12]. The significantly higher melt temperature in die casting of brass, the lower entrance velocity of the melt, the chemical properties of the melt as well as other process parameters clearly separate the casting conditions of brass and aluminium. This motivates the present study, which aims at elucidating the life-limiting failure mechanisms in field-tested brass die casting dies. Two cavity inserts and eight cores were examined. They were made of two hot work tool steels, quenched and tempered to different conditions, and evaluated with respect to hardness changes, growth and density of cracks, and other failure mechanisms.
2.1. Materials
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The two cavity inserts were made of a hot work tool steel, QRO 90 Supreme (Uddeholm Tooling Designation), with the nominal chemical composition (wt.%) of 0.38 C, 0.39 Si, 0.75 Mn, 2.6 Cr, 2.25 Mo, 0.9 V, and Fe balance. The eight cores were made of two different hot work tool steels, QRO 90 Supreme and Hotvar (Uddeholm Tooling Designation). The nominal chemical composition (wt.%) of Hotvar is 0.55 C, 1.0 Si, 0.75 Mn, 2.6 Cr, 2.25 Mo, 0.85 V, and Fe balance. The cavity inserts and cores were hardened and tempered, according to Table 1. The brass alloy, CuZn33Pb2Si–C (Ametal C, Tour and Andersson Designation), with an approximate chemical composition (wt.%), as obtained on a casting by X-ray spectroscopy of 64.1 Cu, 32.3 Zn, 1.9 Pb, 0.71 Si, 0.35 Fe, 0.05 As, 0.03 Al, was used as casting material throughout the field tests. The alloy has a nominal liquidus temperature of 887 ◦ C and a solidus temperature of 844 ◦ C. 2.2. Die casting conditions for the investigated tools All cavity inserts and cores have been used for production of brass valves. The geometry of the studied die components is shown in Fig. 1. The field tests were made in a 2.8 MN cold chamber machine. The temperature of the brass melt was 970 ◦ C and the total cycle time for one casting was 48 s. The entrance velocity of the melt was 1–2 m/s and the maximum metal pressure during injection was approximately 190 MPa. The total shot weight of each casting was 2.1 kg. Water at 20–50 ◦ C was continuously circulated through cooling channels in the cavity insert. After ejection of each casting, the surfaces of the tools were lubricated but not intentionally cooled by spraying a commercial lubricant fluid (Klüber Metalstar Fe 143). The cavity inserts and cores were shot-blasted after each 600 and 200 castings, respectively, and also after the tests. This treatment was performed to remove deposits of Zn compounds on the tool surface and to decrease the tendency of soldering. Note that the cavity inserts in this study were not preheated. The two studied cavity inserts and two of the cores had been used in production until the surface finish or the dimensional changes, or both, of the castings exceeded the tolerances. The tool lifetime of the studied cavity inserts Table 1 Nominal heat treatment and hardness of the field-tested cavity inserts and cores Die component
Steel grade
Austenitising (◦ C/min)
Tempering (◦ C/h)
Cavity insert Cavity insert Core Core
QRO 90 QRO 90 QRO 90 Hotvar
1020/35 1020/34 1020/57 1050/17
640/2 615/2 625/2 580/2
× × × ×
2 2 2 2
Hardness (HV) 430–440 480–500 460–480 600–620
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Fig. 1. The field-tested tool components used in brass die casting. (a) Cavity insert: the arrow indicates the gate where the melt enters. (b) The studied core, located closest to the gate. A and B indicate the positions of the detailed evaluation of the cavity inserts and cores, respectively. The diameter of the cavity at A is 33 mm and that of the core at B is 29 mm.
was 45 000–50 000 castings, whereas that of the cores was 10 000–15 000 castings. Of the eight studied cores, one of each material was investigated after their expired lifetime, whereas one core of each material was taken out of production for examination after 500, 1000, and 1500 castings, respectively. 2.3. Characterisation Identification of the failure mechanisms was performed by macroscopic examinations as well as by detailed studies on the surface of the tools, and on fractured and polished cross-sections, using light optical microscopy (LOM), scanning electron microscopy (SEM), and energy dispersive X-ray spectroscopy (EDS). The fractured cross-sections were obtained after having cooled the steel in liquid nitrogen. For the cavity inserts and cores, the detailed evaluations was focused to an area which was expected to be representative of thermal fatigue cracking with a minimum of stress concentrations (see Fig. 1). The core located closest to the gate was examined since the casting conditions are expected to be most severe where the melt enters the cav-
ity (see Fig. 1b). For example, the melt temperature is at its highest level, the time of contact between the casting alloy and the tool is the longest, and the erosive and corrosive damages resulting from the incoming melt is expected to be most distinguished on that core. Hardness versus depth profiles to reveal any thermal softening were obtained by Vickers indentations on polished cross-sections, using a load of 25 g. The evaluation of the cavity inserts and cores, respectively, were performed at the same location in each die component to minimise any dependence on differences in the casting conditions in the tool.
3. Results 3.1. Macroscopic failure mechanisms in field-tested brass die casting dies The macroscopic surface damages on the two worn-out cavity inserts and cores, respectively, were severe and observed as a network of surface cracks and as individual and clearly pronounced cracks at sharp corners (see Fig. 2). Note
Fig. 2. Typical macroscopic surface damages observed on worn-out cavity inserts and cores. (a) Overview of an insert revealing local severe surface cracking (A) and individual cracks at sharp geometrical corners (B). (b) Overview of a core revealing a severe network of surface cracks, individual cracks at sharp radiuses, and remnants of solidified brass within the cracks (indicated by the arrow).
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Fig. 3. Macroscopic damages observed on cores. (a) Overview after 500 castings revealing individual and clearly pronounced cracks at sharp edge (indicated by the arrow). (b) Overview of a worn-out core revealing a network of surface cracks, and a gross fracture (indicated by the arrow).
the locally severe surface cracking on the worn-out insert, where a protuberance in the tool has been contacted by the melt from several directions. On one of the worn-out cavity inserts, it was observed that some cracks had grown together and caused loss of small steel fragments. In general, the crack networks on the worn-out inserts are not as rough as those on the worn-out cores. Macroscopic inspections revealed local filling of solidified brass in the cracks on the worn-out cores. Typically, individual and clearly developed cracks as well as local filling of solidified brass in the cracks were detected on the cores already after 500 castings (see Fig. 3a). The two worn-out cores showed these defects in a further pronounced stage, in addition to a relatively coarse crack network (see Fig. 3b). Gross fracture was observed on one of the worn-out cores, evident by the missing tip in Fig. 3b. The network of cracks on the two worn-out cores consisted of both wide and narrow cracks, locally filled with solidified brass (see Fig. 4). Note that wide cracks form the network, while more narrow cracks seem to grow into or within the areas separated by the wide cracks (also seen in Fig. 2b). The tendency of brass filling the cracks seems to
Fig. 4. Typical surface damages observed on worn-out cores (SEM). A severe network of wide and narrow cracks and remnants of brass within the cracks are revealed. The vertical scratches on both the surface of the tool (A) and on the solidified brass (B) were made after tool disassemblement.
be larger for the wide than for the narrow cracks. No brass is adhered to the areas separated by the cracks. 3.2. Thermal fatigue cracks 3.2.1. Fractured cross-sections It proved possible to reveal the geometrical details of the cracks by making fractured cross-sections, since the surface of a thermal fatigue crack was easily distinguished from the imposed fracture. At an early stage, the cracks are semi-circular (see Fig. 5). Later on, the crack may grow to a considerable depth, see Fig. 6a, and also extend in the lateral direction (see Fig. 7). The position of the crack front prior to the imposed fracture coincides with a strong concentration of Pb as detected by EDS. Lead also shows up brighter in the micrographs due to atomic number contrast (see Fig. 6a and b). At the crack front of Fig. 6a, a small semi-elliptic sub-crack was detected (see Fig. 6b). A dimple pattern is revealed along the crack tip with bright contrast (see Fig. 6c). 3.2.2. Polished cross-sections Studies of polished cross-sections further elucidate the morphology of the crack pattern (see Fig. 8). The cracks tend to grow perpendicular to the tool surface and reach several
Fig. 5. Fractured cross-section of a core after 500 castings revealing a large semi-circular crack (SEM). The original fatigue crack front is indicated by the arrow.
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Fig. 6. Fractured cross-sections of worn-out cores (SEM). (a) Overview revealing a major crack and a small sub-crack initiated at the front of the major crack (indicated by the arrow). The frame indicate EDS-mapped area, see below. (b) Close-up of the front of the sub-crack of (a). (c) A further increase in magnification reveals, through the characteristic dimples, the presence of a ductile material at the crack tip. (d) EDS map of Pb of the sub-crack of (a). High concentration is represented by bright areas.
millimetres in depth (see Fig. 8a). All of them are filled with a material that has a layered structure next to the steel walls, and a more inhomogeneous structure in the centre where it contains voids and cracks (see Fig. 8b). The appearance of a typical crack tip and the material filling it is seen in Fig. 9. The only element detected at the very crack tip was Pb (see Fig. 9b). Lead also reveals itself in Fig. 9a due to its high atomic number contrast. Most of
Fig. 7. Fractured cross-section of worn-out core revealing a major fatigue crack with the crack front (indicated by the arrow) extending parallel to the surface of the tool (SEM).
the material filling the crack also contains O, indicating a strong element of oxidation (cf. Fig. 9d). The cracks are typically filled already at an early stage of development, and especially after further development (see Fig. 10). Figs. 11 and 12 reveal that their surfaces mainly consist of iron oxides and that the interior is filled with residuals from solidified brass. Pb was not detected in the tip area of this crack. Relatively short cracks (cf. Fig. 13), filled with a mixture of iron and zinc oxides, and a diffuse Cr- and O-rich layer in between the oxides filling the crack and the tool steel, were also frequently observed. Neither Cu nor Pb was detected in these cracks. In general, the appearance of polished cross-sections of the cavity inserts resembles those of the cores. 3.2.3. Crack length and density The maximum length, mean length, and density of the thermal fatigue cracks extracted from polished cross-sections of the tools are of the order of 0.4–3.2 mm, 0.03–1.5 mm, and 0.5–5.5 mm−1 , respectively, for all die components at all number of castings produced (see Fig. 14). After the initial 500 castings, the maximum and mean crack length in the cores is shorter in the harder than in the softer tool, whereas the crack density is higher in the harder tool
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Fig. 8. Polished cross-section of a worn-out core (SEM). (a) Overview revealing typical thermal fatigue cracks filled with material. (b) Close-up of the crack indicated by the arrow in (a).
steel (see Fig. 14). Evidently, the difference in maximum and mean crack length and crack density in the cores averages with number of castings, and reach approximately the same level when they are worn-out, independent of steel grade and initial condition. For the worn-out cavity inserts, the maximum crack length is somewhat shorter and the crack density is significantly lower in the harder than in the softer tool, whereas the mean crack length shows the inverse correlation with hardness. Finally, it is clearly seen that the maximum as well as the mean crack length is significantly shorter in the worn-out cavity inserts than those in the worn-out cores. Evidently, the crack density in
the worn-out die components is higher in the cavity inserts than that in the cores. 3.3. Hardness profiles As expected, for all die components the surface hardness is decreasing with the number of castings produced (see Fig. 15). After the first 500 castings, the hardness reduction of the cores is limited to a surface layer of about 2 mm, and the loss of hardness is larger for the harder than for the softer tool (see Fig. 15a and b). Evidently, the difference in surface hardness between the cores is reduced and the
Fig. 9. Polished cross-section through a typical crack tip in high magnification. (a) Compositional SEM image illustrating the layered filling. (b–d) EDS maps. The element concentration is proportional to the brightness.
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(cf. Fig. 6a, b, and d). Obviously, Pb from the brass casting alloy concentrates at the crack front during casting and due to interaction with the tool steel it is able to locally assist the crack growth. Since the thermal cracking failure dominates, further consideration on this mechanism is given below. 4.1. Thermal cracking
Fig. 10. Polished cross-section of a worn-out core revealing typical long and short cracks (LOM). The frames indicate EDS-mapped areas, see below.
softened surface layer increases in thickness with number of castings. When the cores are worn-out, the softened surface layer extends beyond 10 mm depth. For the worn-out cavity inserts, it is observed that the hardness in the superficial surface layer is reduced to approximately the same level, whereas an approximately constant difference in hardness between the two tools is maintained deeper below the surface (see Fig. 15c). The softened surface layer of the two cavity inserts is limited to a depth of less than 1.5 mm. 4. Discussion It was observed that the dominating failure mechanism in the investigated brass die casting cavity inserts and cores was thermal fatigue cracking (cf. Fig. 2). One of the most striking results from this investigation was that the thermal fatigue crack front advancement is associated with a local enrichment in Pb, whereas it is pinned in areas without Pb
4.1.1. Crack nucleation and growth Thermal cracks are nucleated when the local tensile stress in the surface exceeds a critical level. The tensile stresses are, in turn, a result of plastic yield of the tool material during the hot phase of the casting cycles. It was found that it takes less than 500 brass die casting cycles to nucleate and propagate the first thermal fatigue cracks in the cores (cf. Figs. 3a and 14). Consequently, the temperature alternations in the tool surface layer are severe enough to impose compressive stresses high enough to cause local accumulation of plastic strains to nucleate and propagate thermal cracks in the tool surface by a low-cycle fatigue process. The initial growth of the thermal fatigue cracks is facilitated by an oxidation attack on the crack surfaces to form a Cr- and O-rich layer in between oxides filling the crack and the visually unaffected tool steel (cf. Fig. 13 and Refs. [6,7]). Presence of oxides and brass materials in the cracks increases the compressive stresses and plastic yield during the hot phase of the thermal cycle. The result is a further increase in tensile stress during the cold phase, which results in crack growth. The thermal cracks advance considerably in length while they are typically filled with a mixture of oxides [6], and residuals from the brass alloy [7] (cf. Figs. 10–12). The oxides filling the cracks consist mainly of iron oxides [7], but also of zinc oxides and lead oxides [6], as shown in Figs. 9–13. Their structure in between the steel walls are layered or inhomogeneous and contain cracks
Fig. 11. EDS maps of the area indicated by frame A in Fig. 10. The element concentration is proportional to the brightness.
Fig. 12. EDS maps of the area indicated by frame B in Fig. 10. The element concentration is proportional to the brightness.
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Fig. 13. Polished cross-section through a typical crack of a worn-out core in high magnification: (a) LOM image; (b–f) EDS maps. The element concentration is proportional to the brightness.
and voids [6] (cf. Figs. 8 and 9). The layered structure of the filling material indicates that the crack filling aggravates gradually during the casting process. The tensile stresses imposed on the tool material surface layer during the cooling phase effectively cause local rupture of filled material and open the thermal cracks (cf. Figs. 8b and 13). These cracks in the filling material act as channels for more brass materials to fill the cracks and for oxygen to penetrate down to the crack tip area and oxidise the steel and also the filling material (cf. Figs. 8–12). In addition, the opening of the cracks in the filling material gives an indication of the maximum amount of brass that can fill the crack during the next casting cycle (cf. Fig. 8b). 4.1.2. Crack growth promoted by Pb embrittlement Steel exposed to liquid lead may form solid Fe–liquid Pb systems that causes liquid metal embrittlement [13]. Embrittlement may also occur when the temperature is below the melting point of the embrittling metal (solid metal-induced embrittlement), but the severity of the embrittlement increases considerably at the melting point of the embrittling metal. The mechanism of liquid metal embrittlement is probably associated with a reduction of the atomic bond strength of the solid metal. Previously, temperature recordings in a 5 mm surface layer of a cavity insert during actual die casting of brass revealed that the maximum tool temperature was about 828–450 ◦ C, decreasing from the tool surface and inwards, and that the minimum temperature through the surface layer was about 300 ◦ C [14]. Thus, the temperature on the tool surface layer exceeds the melting point of Pb (327 ◦ C) during each casting cycle and results, consequently, in al-
ternating melting of Pb in the brass alloy filling the cracks (cf. Fig. 10). In this study, it was observed that in a relatively early stage, the thermal fatigue cracks have generated a semi-circular shape (see Fig. 5). Subsequently, the crack front expands in the areas where the imposed stresses are high enough to enable propagation (cf. Fig. 6a). The alternating melting of Pb during each casting cycle cause the liquid Pb from the brass alloy to penetrate through the crack down to the very crack tip, where it is enriched (cf. Figs. 6 and 9b). The enrichment in Pb at the very crack tip is supported by the absence of O in this area (cf. Fig. 9b and d), and also by the observation of a ductile material (Pb) at the crack front (see Fig. 6c). It is obvious that the crack front advances locally where there is an enrichment in Pb, whereas it is pinned in areas without Pb (see Fig. 6a, b, and d). Obviously, the chemical interaction with Pb from the brass alloy and the tool steel is able to locally assist the crack growth, most probably by liquid metal embrittlement. 4.1.3. Crack growth promoted by thermal softening of the tool steel The crack evaluations after the initial 500 castings revealed that the growth of thermal cracks are suppressed by higher tool hardness levels, as evidenced by the shorter maximum and mean crack lengths in the initially harder Hotvar core than in the softer QRO 90 core (see Fig. 14a and b). Thereafter, any differences in the crack characteristics among the cores averages with the number of castings to approximately equal levels at their expired tool life of 10 000–15 000 castings, independent of tool steel and initial condition.
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Fig. 14. Maximum crack length (a), mean crack length (b), and crack density (c) vs. number of castings in the cavity inserts and cores, based on cracks larger than ∼5 m. Each pile is based on measurements along a line of about 17 mm. The designations I and II correspond to cores of QRO 90 (460–480 HV) and Hotvar (600–620 HV), respectively, III and IV to cavity inserts of QRO 90 at nominal hardness values prior to the field tests of 430–440 and 480–500 HV, respectively.
The fact that the average of the crack characteristics with the number of castings (Fig. 14a and b) correlate strongly to the average of the surface layer hardness among them (Fig. 15a and b) indicates that the crack growth is facilitated by the gradual softening of the tool steels. 4.1.4. Crack network development and final appearance Already after 500 castings, individual and clearly developed cracks and crack filling were clearly detected on the cores by the naked eye (cf. Fig. 3a). Subsequently, a network of surface cracks and more individual cracks are formed and the tendency of crack filling aggravates (cf. Figs. 3b and 4). Naturally, the tendency of crack filling is low when the cracks are narrow and increases as they become wider (see Fig. 4). Eventually, some surface cracks may grow together and cause local detachment of surface material (splintering), or deeply penetrate into the tool and lead to gross fracture (cf. Fig. 3b). The observation that the worst surface cracking appeared in tool areas that are exposed to the most critical thermal
Fig. 15. Hardness vs. depth profiles after different number of castings: (a) cores of QRO 90 (460–480 HV); (b) cores of Hotvar (600–620 HV); (c) worn-out cavity inserts of QRO 90.
and/or mechanical stress concentrations, such as sharp corners and protuberances, or parts of the tool that had been more or less surrounded by the melt (cf. Fig. 2), is in agreement with previous findings [5,7]. The fact that the cores are exposed to more intense thermal conditions than the inserts is supported by their more pronounced hardness reduction demonstrated in Fig. 15. Further support is given by the observation that the thermal cracks penetrate considerably deeper into the cores than into the inserts (see Fig. 14a and b). 4.2. Influence of tool maintenance on crack characteristics Previously, it has been reported that the development of thermal fatigue cracks can be characterised by the three consecutive stages: crack nucleation, rapid crack growth, and retarded crack development [5,15,16]. After some initial cycles during which most cracks are formed, some cracks continue to grow deeper into the tool material [5]. Since areas along each thermal crack are locally stress relieved, the growth of adjacent cracks is retarded as the larger once propagate into the tool material [17].
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The tool maintenance applied during the die casting service conditions includes lubrication after each casting and intermittent cleaning by shot-blasting as well as other mechanical procedures, as indicated by the scratches on the tool surface in Fig. 4. From Fig. 14, it is obvious that the crack development in the cores neither can be perfectly characterised by the three consecutive stages, nor by the crack density saturation after the initial castings during which almost all cracks are formed. Surface material removal during tool maintenance is most likely the reason why the crack density in the cores has a tendency to decrease within the whole range of castings produced (see Fig. 14c). Another effect from this maintenance is that the crack length values of Figs. 14a and b are somewhat shorter than they certainly would have been without maintenance. The maximum crack length in the inserts (about 1.3–1.4 mm) after 45 000–50 000 castings (Fig. 14a) is in good agreement with the maximum crack length (about 1.5 mm) estimated from temperature recordings in a cavity insert during actual die casting of brass [14]. The cavity insert, tool steel, brass alloy, and the die casting process parameters during the temperature recordings resemble those of this study. It is therefore concluded that the maintenance has not reduced these maximum crack length values considerably. Finally, the tool maintenance is probably the reason why no macroscopic evidence of erosive or corrosive damage (including soldering) was identified on the investigated tools. However, those modes of tool failure appear to be of secondary importance in the degradation of the investigated tools. It is important to realise that if the field tests had been performed without any tool maintenance, it had not taken many castings before practical casting operation problems arise.
5. Conclusions In this study, the life-limiting failure mechanisms in field-tested cavity inserts and cores for die casting of brass were examined. The following conclusions can be drawn: • The dominating failure mechanism in brass die casting tools is thermal fatigue cracking. • The crack nucleation is associated to accumulation of the local plastic strain in the surface, typical of a low-cycle fatigue process. • The initial growth of the thermal fatigue cracks is facilitated by oxidation of the crack surfaces. • Proceeded growth is facilitated by filling of cracks with brass, by additional oxidation, and by softening of the tool material. • The oxides consist mainly of iron oxides, but also of zinc and lead oxides.
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• The most striking observation is a local enrichment of Pb at the crack front. • Since the crack has advanced locally in these regions, this Pb enrichment is suggested as an important mechanism of crack growth promotion in brass die casting. • In addition to increasing the overall tool steel yield strength and/or oxidation resistance, there is a potential to improve the tool life either by removing or substituting the Pb in the brass casting alloy with any non-harmful element, or using a tool material not susceptible to Pb embrittlement.
Acknowledgements The authors are grateful to Tour and Andersson AB and Mr. Jan Andersson at this company, and to Uddeholm Tooling AB and Bodycote Heat Treatment AB. The financial support from the Swedish Knowledge Foundation is also acknowledged. Special thanks to Dr. Richard Westergård at the Ångström Laboratory for performing the EDS mapping.
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