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Structures of intermetallic phases formed during immersion of H13 tool steel in an Al–11Si–3Cu die casting alloy melt
Materials Science and Engineering A260 (1999) 188 – 196
Structures of intermetallic phases formed during immersion of H13 tool steel in an Al–11Si–3Cu die casting alloy melt Z.W. Chen *, D.T. Fraser 1, M.Z. Jahedi CSIRO Manufacturing Science and Technology, Locked Bag No. 9, Preston, Melbourne, Victoria 3072, Australia Received 1 June 1998; received in revised form 24 August 1998
Abstract The structures of intermetallic alloy layers formed during immersion of H13 tool steel into an aluminium die casting alloy melt have been studied by X-ray diffraction. Energy dispersive spectroscopy (EDS) analysis on the intermetallic phases was also conducted. A thick composite layer away from the H13 steel substrate consisted of irregular intermetallics and solidified cast alloy. A thin intermetallic layer was present between the thick composite layer and an inner compact layer next to the steel substrate. The intermetallic phase in the composite layer was found to have a cubic structure, abcc-(FeSiAlCrMnCu). The thin layer was identified to be structurally isomorphous with hexagonal aH-Fe2SiAl8. The compositional difference between aH and abcc intermetallic phases was mainly that the latter consisted of a higher amount of Cr + Mn +Cu. This is consistent with the suggestion that chromium, manganese and copper stabilise abcc phase at the expense of aH phase. The inner compact layer next to the steel substrate was identified to be isomorphous with orthorhombic h-Fe2Al5. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Intermetallic alloy layers; H13 tool steel; Immersion test
1. Introduction High pressure die casting (HPDC) is a widely used process to produce near net shaped components of aluminium alloys. The process consists of injecting molten alloy into a die cavity at high velocities and then solidification of the alloy at high rates. H13 tool steel is normally used as the die material. Die failure is an important issue in HPDC and one of the major reasons for die failure is soldering [1]. Soldering is the term used in the HPDC industry and it originates from reaction between the die and the casting alloy. As a result of soldering, the solidified alloy can stick to the die, forming defective castings upon ejection. The remains of hard intermetallic phases can build up on the die surfaces and production needs to be stopped to remove * Corresponding author. Tel.: + 61-3-96627850; fax: +61-396627851; e-mail: [email protected]. 1 Present address: Co-operative Research Centre for Alloy and Solidification Technology (CAST): The University of Queensland, Queensland 4072, Australia..
the soldered layer by polishing, which is costly in terms of lost production time, labour cost for polishing, and shortening of die life. Hence, soldering and its prevention has recently been the subject of a considerable research effort [1–8]. In understanding the soldering reaction and the resulting intermetallic phases, immersion tests of H13 tool steel into Al–Si–Cu–Mg alloy melts have been used to study the nature of these phases [3,6]. Sundqvist and Hogmark detected the presence of silicon and chromium in significant amounts in the intermetallic phases using X-ray energy dispersive spectroscopy (EDS) analysis [6]. However, they suggested that the phases are Fe–Al binary phases based on the binary Fe–Al phase diagram. Yu et al. [3] suggested a series of ternary Fe–Al–Si intermetallic phases based on EDS analysis and the Fe–Al–Si ternary diagram adapted by Rivlin and Raynor [9]. However, as shown by Rivlin and Raynor, the phases near the iron corner of the system are far from certain and direct information on equilibria below the liquidus was scanty. The presence
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 0 9 6 3 - 0
Z.W. Chen et al. / Materials Science and Engineering A260 (1999) 188–196 Table 1 Composition of the H13 tool steel (wt%)
189
C
Cr
Si
Mo
V
Mn
Fe
have provided a set of microstructures for phase identification, which are typical of those observed previously.
0.4
5.3
1.0
1.3
0.9
0.4
Bal.
2.2. Sample preparation and analysis
of chromium and other elements in the phases adds further uncertainty. As part of the research effort to understand the metallurgy of soldering in HPDC, the present work was undertaken to identify the major intermetallic phases formed during immersion of H13 tool steel in an Al– 11Si–3Cu melt. The structures of the phases were determined using X-ray diffraction. The identification of the phases was further assisted by compositional analysis using EDS.
2. Experimental procedure
2.1. Materials and immersion test H13 tool steel samples (20×10 ×80 mm) were used for the immersion test and the composition is given in Table 1. A silicon carbide crucible was used in an induction furnace to melt approximately 1 kg of Al– 11Si–3Cu alloy with composition as given in Table 2. The temperature of the melt was 680 9 2°C, measured using a k-type thermocouple. H13 steel samples were polished to 600 grit before immersion into the melt. The duration of immersion was 5 h and samples were quenched in water immediately after withdrawal from the melt. Immersed samples were then prepared for both metallography and X-ray diffraction. The SEM micrograph in Fig. 1 shows the morphologies of the intermetallic phases formed and grown during immersion of a H13 steel sample in the melt. On top of the steel substrate is a thin compact layer (10–15 mm). Next to this, there is another thinner layer, the outer compact layer, before a thick composite layer of irregular morphology. The interface between the outer compact layer and the thick composite layer is highly irregular, as though the irregular phase in the composite layer was broken away from the compact layer. The composite layer is followed by solidified cast alloy. The existence of the compact layer and a thicker composite layer can be found in micrographs given in the previously quoted studies [3,6]. The immersed samples
Successive diffractometer scans were made on the composite layer and the compact layers. Samples were prepared by mechanical thinning using SiC paper (to 1200 mesh) and Fig. 1 shows the approximate locations where representative results have been presented in this paper. Mechanical thinning to the composite layer was easily done as the layer was thick (: 200 mm). For the combined compact layer samples (OC in Fig. 1), any slight angle of the thinning plane to the compact layer plane would mean one side of the sample was being thinned to or near the H13 substrate. Hence, diffraction from the substrate could be expected. As already discussed, the outer compact layer is very thin. A slight angle of the thinning plane and a slight difference in vertical position of machining would mean that the composition of the phases on the OC sample surfaces differed significantly. If most of the outer compact layer was preserved, the sample was likely to contain a significant amount of the composite layer materials. If the composite layer materials were totally machined off, a significant amount of the outer compact layer was also likely to be removed. Hence, only weak diffraction from the outer compact layer was to be expected. It was also expected that diffraction intensities from the inner compact layer would be strong in both OC1 and OC2 patterns. During preparation of samples, sometimes cracking occurred in the inner compact layer near and along the interface between the compact intermetallic layer and H13 steel substrate. A diffraction pattern was obtained from a sample on which the coated materials had cracked and peeled off; in other words, from a H13 steel sample with residual inner compact layer material on the surface. This location is indicated in Fig. 1 as IC-H13 pattern (inner compact and H13), where a slightly irregular surface has been drawn to show that this was a fracture surface rather than a normal ground surface. When diffraction from H13 steel was excluded, the remainder of the diffraction pattern was expected to be purely from the inner compact layer material. Diffraction experiments were performed using a Siemens D500 X-ray diffractometer with monochroma˚ ). The scanning tized CuKa1 radiation (l=1.54056 A
Table 2 Composition of the cast alloy melt (wt%) Al
Si
Cu
Fe
Zn
Mg
Mn
Cr
Ni
Ti, Sn, Sb
Bal.
11.5
2.8
1.0
1.0
0.16
0.20
0.03
0.1
0.03
190
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Fig. 1. Microstructures of H13 tool steel immersed in the Al – 11Si – 3Cu melt at 680°C for 5 h and quenched in water. Approximate locations for producing X-ray diffraction patterns presented in this paper are also illustrated.
Fig. 2. X-ray diffraction pattern of a composite layer sample (see Fig. 1 for location) showing the existence of phases: a-Al, Si, and abcc-(FeSiAlCrMnCu).
range (2u) was normally from 10 to 90° and the scanning rate was 0.2° min − 1 with a step width of 0.04°. The sample was rotated during the diffraction experiment. Data from the Joint Committee on Powder Diffraction Standards [10] (JCPDS) were used to identify the phases. The identification of the abcc-(FeAlSiCrMnCu) phase was based on a match of the observed d-spacings to the calculated d-spacings. The structure of this phase and its lattice parameter will be discussed in detail in Section 3.1. The symbol (FeAlSiCrMnCu) is used to represent the uncertainty of the stoichiometry.
Semi-quantitative compositional analysis of phases was carried out using a Leica S440 scanning electron microscope (at 15 kV) fitted with a Link Ge EDS detector and Link ISIS software. Analytical spots were mainly chosen in the intermetallic phases. Probe current was set at 2.5 nA and a counting time of 100 s was used for each spot analysis. During each analysis, the spectrum was accumulated for later processing and ZAF correction. In analysis of the irregular intermetallic phase in the composite layer, only those particles with a width approximately larger than 2 mm were used.
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191
Table 3 Observed d-spacings for the irregular intermetallic phase in the composite layer and the calculated d-spacings based on a body centre cubic structure (abcc phase) with a modified lattice parameter a=1258.8 pm dobs. (pm)
dcal. (pm)
hkl
dobs. (pm)
dcal. (pm)
hkl
445.1 398.2 363.9 335.6 296.9 281.3 268.0 246.6 229.6 222.4 215.8
445.1 398.1 363.4 336.4 296.7 281.5 268.4 246.9 229.8 222.5 215.9
220 310 222 321 330, 411 420 332 431, 510 521 440 433, 530
209.5 204.0 199.0 194.0 185.4 174.6 150.4 148.2 146.2 127.1 125.7
209.8 204.2 199.0 194.2 185.6 174.6 150.5 148.4 146.3 127.2 125.9
442, 532, 620 541 631 640 653 660, 743, 770, 860
600 611
822 750, 831 853
Fig. 3. X-ray diffraction pattern of the outer compact layer sample 1 (OC1 in Fig. 1) showing the existence of phases: aH-Fe2SiAl8, abcc-(FeSiAlCrMnCu), h-Fe2Al5, and a-Al, Si, and a-Fe.
3. Results and discussion
3.1. Structures of the composite layer Fig. 2 is a representative X-ray diffraction pattern for the composite layer (see Fig. 1 for location). As the layer was a composite of irregular intermetallic phase and the solidified cast alloy, diffraction peaks from a-Al and silicon were clearly identified in this diffraction pattern. Apart from the a-Al and silicon diffraction peaks, there were 22 peaks clearly identifiable. Their observed d-spacings are listed in Table 3. Structures with these d-spacings could not be identified using JCPDS data for Fe – Al binary and Fe – Al – Si ternary phases. Hence, the irregular intermetallic phase is not a binary Fe–Al nor a ternary Fe – Al – Si phase.
It has been established that the presence of chromium, manganese, and copper in Al–Si alloys containing iron transform hexagonal aH-Fe2SiAl8 to a nearly body centred cubic abcc phase with a=1255 pm [11,12]. In the present case, the chromium content (5.3%) was high in the H13 steel and the cast alloy contained manganese (0.2%) and copper (2.8%). Hence, it is highly probable that the irregular intermetallic phase in the thick composite layer was an abcc phase rather than a ternary aH-Fe2SiAl8 phase. Interplaner spacings were calculated based on a bcc (body centred cubic) structure with a modified lattice parameter. The modification was made by minimizing the value of (2ucal. − 2uobs)2. The calculated d-spacings based on this modified lattice parameter (a=1258.8 pm) are also listed in Table 3. It can be seen in Table 3
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Fig. 4. X-ray diffraction pattern of the outer compact layer sample 2 (OC2 in Fig. 1) showing that the aH-Fe2SiAl8 phase is still present but the major phase is h-Fe2Al5. a-Fe diffraction peaks are clearly present.
that there was good agreement between the observed and calculated d-spacings and hence the intermetallic phase in the composite layer was identified to have a cubic structure and named abcc-(FeAlSiCrMnCu).
3.2. Structure of the outer compact layer Two diffraction patterns of the combined compact layer samples, OC1 and OC2, are presented in Figs. 3 and 4, respectively. Many diffraction peaks from abcc phase as well as those from silicon and a-Al can be identified in the OC1 pattern but not in the OC2 pattern, meaning that the OC1 sample surface contained a significant amount of composite layer materials. Diffraction peaks from the inner compact layer can be identified in these two patterns and this will be discussed in detail in Section 3.3. Diffraction peaks from the H13 steel substrate are clearly identified in both the OC1 and OC2 samples. The reason for this has already been explained in Section 2.2. Table 4 lists the diffraction peaks excluding all those identified for the composite layer materials (abcc phase, silicon, a-Al), the inner compact layer and the H13 steel. In Table 4, data from JCPDS file for aH-Fe2SiAl8 are also listed. From Fig. 3 and Table 4, a number of diffraction peaks in OC1 identify positively the outer compact layer as aH-Fe2SiAl8. These are aH 204, 205, 124, 306, 228, 1210 (underlined for 1] 10), and 239. The strongest peaks for the phase were 237 and 330 but they were strongly influenced by the strong peak of abcc 532, 611, and the very strong peaks of a-Al 200 and a-Fe 110. aH 236 (d= 214.1 pm), 501 (d = 213.6 pm), and 3010 (d =211.0 pm) as listed in the JCPDS file could not be used as they
are strongly influenced by abcc 433, 530, and h 002. Not all the diffraction peaks listed in the JCPDS file for aH phase could be detected in pattern OC1 (Fig. 3). Each JCPDS data file is based on a diffraction experiment using a sufficient amount of powder. In the present case the outer compact layer is very thin, as shown in Fig. 1. Hence, it was reasonable for many weak diffraction peaks as listed in JCPDS to be absent. On the other hand, JCPDS for the aH phase only lists diffraction peaks with d-spacings \ 170 pm (2u = 54°). This is why the weak but distinctive peak at 2u= 73.54° in pattern OC1 has been labeled as UI (unidentified). For aH 3511 (d=128.7 pm and 2u= 73.52°) and aH 457 (d= 128.8 pm and 2u= 73.53°), the multiplicity factor is high (=24). For a hexagonal structure, the structure factor for 3511 and 457 is also high ( = 0.75f 2 where f is the scattering factor). Hence, it was likely that the peak at 2u= 73.54° was due to aH 3511, 457. In pattern OC2 (Fig. 4 and Table 4), diffraction peaks from the abcc phase, silicon and a-Al were absent meaning that the surface of the sample contained very little composite layer materials. This also suggests, as already discussed in Section 2.2, that only a very small amount of the outer compact layer was left on the sample surface and hence the intensities from aH diffraction were expected to be weak. Peaks from aH 112, 204, 205, 124, 137, 239 could be detected in the pattern. Due to the absence of abcc 532, 611 in pattern OC2, aH 237, which is one of the strongest diffraction peaks for the phase, was clearly identified. Furthermore, the aH 330 peak could also be identified.
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193
Table 4 Observed and calculated a (a= 1237.2 pm, c = 2615.8 pm) d-spacings, X-ray diffraction intensities for the phase isomorphous with aH-Fe2SiAl8 in patterns OC1 and OC2 and from JCPDS dobs (pm)b
dcal a(pm)
IOC1c
IOC2c
IJCPDS
558.9
559.2 535.7 524.7 495.8 456.5 449.4 436.0 414.4 400.2 386.9 374.3 367.3 357.2 352.8 344.3 330.5 337.0 320.2 306.5 296.7 295.0 289.8 289.1 279.6 276.3 274.6 270.5 267.9 255.5 254.4 252.3 245.6 244.7 241.2 238.4 236.6 233.8 232.6 228.2
vw
vw
14 8 16 7 10 2 8 12 44 5 12 6 16 19 34 16 2 28 1 4 5 2 2 3 5 1 2 2 2 7 3 8 3 2 4 10 36 18 9
414.3 399.9 374.3
344.4
320.2 296.5
276.2
245.8
232.4
w —d
vw vw
w
vw
m
mw
—e
—e
—f
vw
—g
mh
w
hkl 112 200 201 202 203 114 006 204 121 122 205 123 300 107 124 303 008 125 207 216 305 132 118 224 306 127 134 400 209 128 226 136 231 308 405 233 140 137 406
dobs (pm)
224.9 222.3 219.7
206.4 205.3
199.5
187.5
185.9
177.6
dcala (pm) 225.4 224.7 222.4 219.7 218.0 217.4 214.1 213.6 211.6 211.0 208.1 207.2 206.2 205.4 203.6 201.9 199.7 198.3 197.9 197.2 192.3 190.4 188.4 187.7 186.8 186.4 186.1 185.7 184.6 184.1 180.6 178.6 177.8 176.9 176.2 171.6 170.1 169.7
IOC1c
IOC2c
w —i w
—j —k —k —k
—k
—l sm —n
mwl mm —n
mwo
m
w
mwp
vw
w
IJCPDS
hkl
4 4 13 6 26 39 44 31 7 80 4 11 100 98 49 3 12 30 27 7 5 9 13 28 4 8 12 13 3 2 2 2 7 2 1 1 2 2
309 228 235 1210 0012 2011 236 501 502 3010 503 408 330 237 504 2012 2210 505 3011 243 506 152 2013 239 0014 336 3012 1311 154 1014 155 600 4011 602 430 520 434 2114
a
Based on the observed d-spacings as bolded. d-spacings bolded identify the phase with high confidence. c Intensity: vs, very strong; s, strong; m, medium; mw, medium to weak; w, weak; vw, very weak. d Overlaps with medium-strong abcc 310 peak. e Overlaps with medium-strong h 111 peak. f Overlaps with weak abcc 411 330 peak. g Overlaps with weak abcc 431 510. h Overlaps with weak Al 111. i Overlaps with weak abcc 440 peak. j Overlaps with medium-strong abcc 433 530 peak. k Overlaps with very strong h 002 peak. l Overlaps with strong aH 237 peak and very strong peak of combined Al 200 and Fe 110 in OC1 but more distinctively identified in OC2. m Overlap with very strong Al 220 and Fe 110 peak in OC1 but not in OC2. n Overlaps with strong abcc 532 611 peak in OC1 and strong Fe 110 in OC2. o Overlaps with weak abcc 620 peak. p Overlaps with weak abcc 631 peak. b
Richards et al. [13] recently used X-ray diffraction to identify the structures of intermetallic phases present in aluminized coatings. They identified, for hot dipping of
plain carbon steels in an Al–11Si melt, the outer intermetallic layer overlying the h-Fe2Al5(Si) layer to be aH-Fe2SiAl8. This is consistent with our identification of
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Table 5 Observed and calculated (a= 766.2 pm, b = 637.5 pm, and c =425.7 pm) d-spacings, X-ray diffraction intensities for the phase isomorphous with h–Fe2Al5 present in layers patterns OC2 and IC-H13 dobs. (pm)
dcal. (pm)
hkl
IOC2
IIC-H13
321.4 213.0 195.2 177.0 147.6 136.1 122.7 121.8
321.4 212.9 195.2 177.0 147.6 136.3 122.8 121.8
111 002 112 022 132 113 223 313
m vs m w vw w mw mw
mw s mw – – – – –
the outer compact layer to be aH phase, although detailed comparison could not be made, as Richards et al. did not give their detailed diffraction patterns. The abcc phase has not been reported to be present in aluminized coatings [13]. aH-Fe2SiAl8 is a stable Fe– Al–Si ternary phase [14] and the absence of the abcc phase is to be expected when aluminizing of plain carbon steels in an Al – Si – (Fe) melt.
3.3. Structure of the inner compact layer Strong diffraction from the inner compact layer was to be expected in patterns OC1 and OC2, as already explained in Section 2.2. As discussed, diffraction from abcc, silicon, a-Al of the residual composite layer and aH from the outer compact layer are found in pattern OC1 while only diffraction peaks of medium to weak
intensity from aH phase are present in pattern OC2. Hence, identification of the inner compact layer using pattern OC2 (Fig. 4) was simpler and clearer. Table 5 lists the data for pattern OC2 excluding those diffraction peaks due to the aH phase and the H13 substrate. From Fig. 4 and Table 5, eight diffraction peaks were identified for the structure isomorphous with orthorhombic h-Fe2Al5. This structural identification of the inner compact layer is consistent with that of the layer formed next to the steel substrate during aluminizing using a Al–11Si melt as a phase structurally isomorphous with h-Fe2Al5 [13]. Diffraction from the h-phase in patterns OC1 (Fig. 3) and OC2 (Fig. 4) are similar. All the h-Fe2Al5 diffraction peaks listed in Table 5 can be identified in pattern OC1. Further positive identification of the inner compact layer being a phase isomorphous with the h-Fe2Al5 was provided by the IC-H13 sample (Fig. 5 and Table 5). This diffraction pattern shows 111, 002, and 112 from the h-phase which was the residual material (of the inner compact layer) on the H13 immersed sample. From pattern IC-H13, where no diffraction peaks from the aH phase are present, the intensity of 111 is comparable to that of 112. For this reason, it can be concluded that in pattern OC2 (Fig. 4), the peak at 2u= 27.81° whose intensity is also comparable to h 112 in that pattern, was mainly contributed by h 111. Figs. 3–5 and Table 5 show that 002 from the h-phase had a very strong intensity, indicating that the phase has a preferred growth direction such that (001) of the phase was parallel to the steel substrate. This is
Fig. 5. X-ray diffraction pattern (scanning started from 20°) of the H13 sample (IC-H13 in Fig. 1) showing some amount of h-Fe2Al5 phase present on the sample surface.
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Fig. 6. EDS analysis of the intermetallic phases plotted as a function of the distance from the h – aH interface. Selected elements from the analyses on H13 are also plotted.
in common with the preferred orientation of the hFe2Al5 phase found in hot dipping of steel in aluminium melts or melts containing aluminium. The phase grows in the 001 direction towards the steel substrate in a variety of dipping conditions [13,15,16].
3.4. Compositional features of the intermetallic phases Fig. 6 shows a summarized EDS analysis on the intermetallic phases. For the h-phase layer, the average composition for major elements was 51% Al, 42% Fe, 4.5% Si, and 2.5% Cr. This compares to :52%Al, 45%Fe, and 3%Si for the h-phase layer formed during hot dipping of mild steel in an Al – 8%Si melt as recently given by El-Mahallawy et al. [17]. It seems that, in the present case, chromium may replace some iron in the h-phase. In El-Mahallawy et al. samples, precipitates rich in silicon (\ 10%) were clearly present in their h-phase layer. It is not known if these precipitates formed during hot dipping or during the slow cooling (in air) of the samples. Precipitation of the silicon-rich phase in the h-phase layer may have resulted in a decrease in silicon content in the h-phase in Mahallawy et al. samples. From the analytical data presented in Fig. 6, it appears that the chromium content (: 1.5%) was generally lower in the thin aH phase layer than that of the inner h-phase layer. It reverses to a higher content at
approximately 3% in the abcc phase in the composite layer and the iron content decreased slightly from aH to abcc phase. Manganese and copper contents were also lower in the aH phase than in the abcc phase. These compositional changes are consistent with the established phenomenon that chromium, manganese, and copper are the elements required to stabilise a cubic structure (abcc phase) at the expense of the hexagonal structure (aH-Fe2SiAl8) [9,12]. 4. Conclusion The three reaction layers formed during immersion of H13 tool steel in a die casting alloy melt have been structurally identified. The irregular intermetallic phase in the thick composite layer away from the H13 steel substrate was identified to have a bcc structure, abcc(FeSiAlCrMnCu). The thin and continuous layer between the composite layer and the inner compact layer was found to be structurally isomorphous with aHFe2SiAl8. The compositional difference between aH and abcc phases was mainly that the latter consisted of a higher amount of Cr+ Mn+Cu and a lower amount of iron. It is likely that the presence of chromium, manganese, and copper in the system resulted in the transformation aH abcc. The inner compact layer next to the steel substrate was identified to be orthorhombic h-Fe2Al5 containing silicon and chromium.
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Acknowledgements We would like to thank the Co-operative Research Centre for Alloy and Solidification Technology and Nissan Casting Australia for supporting this work. We would also like to thank N.M. Rockelmann and Dr D.G. Hay for the production of X-ray diffraction patterns.
[6] [7] [8] [9] [10]
[11]
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Structures of intermetallic phases formed during immersion of H13 tool steel in an Al–11Si–3Cu die casting alloy melt Z.W. Chen *, D.T. Fraser 1, M.Z. Jahedi CSIRO Manufacturing Science and Technology, Locked Bag No. 9, Preston, Melbourne, Victoria 3072, Australia Received 1 June 1998; received in revised form 24 August 1998
Abstract The structures of intermetallic alloy layers formed during immersion of H13 tool steel into an aluminium die casting alloy melt have been studied by X-ray diffraction. Energy dispersive spectroscopy (EDS) analysis on the intermetallic phases was also conducted. A thick composite layer away from the H13 steel substrate consisted of irregular intermetallics and solidified cast alloy. A thin intermetallic layer was present between the thick composite layer and an inner compact layer next to the steel substrate. The intermetallic phase in the composite layer was found to have a cubic structure, abcc-(FeSiAlCrMnCu). The thin layer was identified to be structurally isomorphous with hexagonal aH-Fe2SiAl8. The compositional difference between aH and abcc intermetallic phases was mainly that the latter consisted of a higher amount of Cr + Mn +Cu. This is consistent with the suggestion that chromium, manganese and copper stabilise abcc phase at the expense of aH phase. The inner compact layer next to the steel substrate was identified to be isomorphous with orthorhombic h-Fe2Al5. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Intermetallic alloy layers; H13 tool steel; Immersion test
1. Introduction High pressure die casting (HPDC) is a widely used process to produce near net shaped components of aluminium alloys. The process consists of injecting molten alloy into a die cavity at high velocities and then solidification of the alloy at high rates. H13 tool steel is normally used as the die material. Die failure is an important issue in HPDC and one of the major reasons for die failure is soldering [1]. Soldering is the term used in the HPDC industry and it originates from reaction between the die and the casting alloy. As a result of soldering, the solidified alloy can stick to the die, forming defective castings upon ejection. The remains of hard intermetallic phases can build up on the die surfaces and production needs to be stopped to remove * Corresponding author. Tel.: + 61-3-96627850; fax: +61-396627851; e-mail: [email protected]. 1 Present address: Co-operative Research Centre for Alloy and Solidification Technology (CAST): The University of Queensland, Queensland 4072, Australia..
the soldered layer by polishing, which is costly in terms of lost production time, labour cost for polishing, and shortening of die life. Hence, soldering and its prevention has recently been the subject of a considerable research effort [1–8]. In understanding the soldering reaction and the resulting intermetallic phases, immersion tests of H13 tool steel into Al–Si–Cu–Mg alloy melts have been used to study the nature of these phases [3,6]. Sundqvist and Hogmark detected the presence of silicon and chromium in significant amounts in the intermetallic phases using X-ray energy dispersive spectroscopy (EDS) analysis [6]. However, they suggested that the phases are Fe–Al binary phases based on the binary Fe–Al phase diagram. Yu et al. [3] suggested a series of ternary Fe–Al–Si intermetallic phases based on EDS analysis and the Fe–Al–Si ternary diagram adapted by Rivlin and Raynor [9]. However, as shown by Rivlin and Raynor, the phases near the iron corner of the system are far from certain and direct information on equilibria below the liquidus was scanty. The presence
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 0 9 6 3 - 0
Z.W. Chen et al. / Materials Science and Engineering A260 (1999) 188–196 Table 1 Composition of the H13 tool steel (wt%)
189
C
Cr
Si
Mo
V
Mn
Fe
have provided a set of microstructures for phase identification, which are typical of those observed previously.
0.4
5.3
1.0
1.3
0.9
0.4
Bal.
2.2. Sample preparation and analysis
of chromium and other elements in the phases adds further uncertainty. As part of the research effort to understand the metallurgy of soldering in HPDC, the present work was undertaken to identify the major intermetallic phases formed during immersion of H13 tool steel in an Al– 11Si–3Cu melt. The structures of the phases were determined using X-ray diffraction. The identification of the phases was further assisted by compositional analysis using EDS.
2. Experimental procedure
2.1. Materials and immersion test H13 tool steel samples (20×10 ×80 mm) were used for the immersion test and the composition is given in Table 1. A silicon carbide crucible was used in an induction furnace to melt approximately 1 kg of Al– 11Si–3Cu alloy with composition as given in Table 2. The temperature of the melt was 680 9 2°C, measured using a k-type thermocouple. H13 steel samples were polished to 600 grit before immersion into the melt. The duration of immersion was 5 h and samples were quenched in water immediately after withdrawal from the melt. Immersed samples were then prepared for both metallography and X-ray diffraction. The SEM micrograph in Fig. 1 shows the morphologies of the intermetallic phases formed and grown during immersion of a H13 steel sample in the melt. On top of the steel substrate is a thin compact layer (10–15 mm). Next to this, there is another thinner layer, the outer compact layer, before a thick composite layer of irregular morphology. The interface between the outer compact layer and the thick composite layer is highly irregular, as though the irregular phase in the composite layer was broken away from the compact layer. The composite layer is followed by solidified cast alloy. The existence of the compact layer and a thicker composite layer can be found in micrographs given in the previously quoted studies [3,6]. The immersed samples
Successive diffractometer scans were made on the composite layer and the compact layers. Samples were prepared by mechanical thinning using SiC paper (to 1200 mesh) and Fig. 1 shows the approximate locations where representative results have been presented in this paper. Mechanical thinning to the composite layer was easily done as the layer was thick (: 200 mm). For the combined compact layer samples (OC in Fig. 1), any slight angle of the thinning plane to the compact layer plane would mean one side of the sample was being thinned to or near the H13 substrate. Hence, diffraction from the substrate could be expected. As already discussed, the outer compact layer is very thin. A slight angle of the thinning plane and a slight difference in vertical position of machining would mean that the composition of the phases on the OC sample surfaces differed significantly. If most of the outer compact layer was preserved, the sample was likely to contain a significant amount of the composite layer materials. If the composite layer materials were totally machined off, a significant amount of the outer compact layer was also likely to be removed. Hence, only weak diffraction from the outer compact layer was to be expected. It was also expected that diffraction intensities from the inner compact layer would be strong in both OC1 and OC2 patterns. During preparation of samples, sometimes cracking occurred in the inner compact layer near and along the interface between the compact intermetallic layer and H13 steel substrate. A diffraction pattern was obtained from a sample on which the coated materials had cracked and peeled off; in other words, from a H13 steel sample with residual inner compact layer material on the surface. This location is indicated in Fig. 1 as IC-H13 pattern (inner compact and H13), where a slightly irregular surface has been drawn to show that this was a fracture surface rather than a normal ground surface. When diffraction from H13 steel was excluded, the remainder of the diffraction pattern was expected to be purely from the inner compact layer material. Diffraction experiments were performed using a Siemens D500 X-ray diffractometer with monochroma˚ ). The scanning tized CuKa1 radiation (l=1.54056 A
Table 2 Composition of the cast alloy melt (wt%) Al
Si
Cu
Fe
Zn
Mg
Mn
Cr
Ni
Ti, Sn, Sb
Bal.
11.5
2.8
1.0
1.0
0.16
0.20
0.03
0.1
0.03
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Fig. 1. Microstructures of H13 tool steel immersed in the Al – 11Si – 3Cu melt at 680°C for 5 h and quenched in water. Approximate locations for producing X-ray diffraction patterns presented in this paper are also illustrated.
Fig. 2. X-ray diffraction pattern of a composite layer sample (see Fig. 1 for location) showing the existence of phases: a-Al, Si, and abcc-(FeSiAlCrMnCu).
range (2u) was normally from 10 to 90° and the scanning rate was 0.2° min − 1 with a step width of 0.04°. The sample was rotated during the diffraction experiment. Data from the Joint Committee on Powder Diffraction Standards [10] (JCPDS) were used to identify the phases. The identification of the abcc-(FeAlSiCrMnCu) phase was based on a match of the observed d-spacings to the calculated d-spacings. The structure of this phase and its lattice parameter will be discussed in detail in Section 3.1. The symbol (FeAlSiCrMnCu) is used to represent the uncertainty of the stoichiometry.
Semi-quantitative compositional analysis of phases was carried out using a Leica S440 scanning electron microscope (at 15 kV) fitted with a Link Ge EDS detector and Link ISIS software. Analytical spots were mainly chosen in the intermetallic phases. Probe current was set at 2.5 nA and a counting time of 100 s was used for each spot analysis. During each analysis, the spectrum was accumulated for later processing and ZAF correction. In analysis of the irregular intermetallic phase in the composite layer, only those particles with a width approximately larger than 2 mm were used.
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191
Table 3 Observed d-spacings for the irregular intermetallic phase in the composite layer and the calculated d-spacings based on a body centre cubic structure (abcc phase) with a modified lattice parameter a=1258.8 pm dobs. (pm)
dcal. (pm)
hkl
dobs. (pm)
dcal. (pm)
hkl
445.1 398.2 363.9 335.6 296.9 281.3 268.0 246.6 229.6 222.4 215.8
445.1 398.1 363.4 336.4 296.7 281.5 268.4 246.9 229.8 222.5 215.9
220 310 222 321 330, 411 420 332 431, 510 521 440 433, 530
209.5 204.0 199.0 194.0 185.4 174.6 150.4 148.2 146.2 127.1 125.7
209.8 204.2 199.0 194.2 185.6 174.6 150.5 148.4 146.3 127.2 125.9
442, 532, 620 541 631 640 653 660, 743, 770, 860
600 611
822 750, 831 853
Fig. 3. X-ray diffraction pattern of the outer compact layer sample 1 (OC1 in Fig. 1) showing the existence of phases: aH-Fe2SiAl8, abcc-(FeSiAlCrMnCu), h-Fe2Al5, and a-Al, Si, and a-Fe.
3. Results and discussion
3.1. Structures of the composite layer Fig. 2 is a representative X-ray diffraction pattern for the composite layer (see Fig. 1 for location). As the layer was a composite of irregular intermetallic phase and the solidified cast alloy, diffraction peaks from a-Al and silicon were clearly identified in this diffraction pattern. Apart from the a-Al and silicon diffraction peaks, there were 22 peaks clearly identifiable. Their observed d-spacings are listed in Table 3. Structures with these d-spacings could not be identified using JCPDS data for Fe – Al binary and Fe – Al – Si ternary phases. Hence, the irregular intermetallic phase is not a binary Fe–Al nor a ternary Fe – Al – Si phase.
It has been established that the presence of chromium, manganese, and copper in Al–Si alloys containing iron transform hexagonal aH-Fe2SiAl8 to a nearly body centred cubic abcc phase with a=1255 pm [11,12]. In the present case, the chromium content (5.3%) was high in the H13 steel and the cast alloy contained manganese (0.2%) and copper (2.8%). Hence, it is highly probable that the irregular intermetallic phase in the thick composite layer was an abcc phase rather than a ternary aH-Fe2SiAl8 phase. Interplaner spacings were calculated based on a bcc (body centred cubic) structure with a modified lattice parameter. The modification was made by minimizing the value of (2ucal. − 2uobs)2. The calculated d-spacings based on this modified lattice parameter (a=1258.8 pm) are also listed in Table 3. It can be seen in Table 3
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Fig. 4. X-ray diffraction pattern of the outer compact layer sample 2 (OC2 in Fig. 1) showing that the aH-Fe2SiAl8 phase is still present but the major phase is h-Fe2Al5. a-Fe diffraction peaks are clearly present.
that there was good agreement between the observed and calculated d-spacings and hence the intermetallic phase in the composite layer was identified to have a cubic structure and named abcc-(FeAlSiCrMnCu).
3.2. Structure of the outer compact layer Two diffraction patterns of the combined compact layer samples, OC1 and OC2, are presented in Figs. 3 and 4, respectively. Many diffraction peaks from abcc phase as well as those from silicon and a-Al can be identified in the OC1 pattern but not in the OC2 pattern, meaning that the OC1 sample surface contained a significant amount of composite layer materials. Diffraction peaks from the inner compact layer can be identified in these two patterns and this will be discussed in detail in Section 3.3. Diffraction peaks from the H13 steel substrate are clearly identified in both the OC1 and OC2 samples. The reason for this has already been explained in Section 2.2. Table 4 lists the diffraction peaks excluding all those identified for the composite layer materials (abcc phase, silicon, a-Al), the inner compact layer and the H13 steel. In Table 4, data from JCPDS file for aH-Fe2SiAl8 are also listed. From Fig. 3 and Table 4, a number of diffraction peaks in OC1 identify positively the outer compact layer as aH-Fe2SiAl8. These are aH 204, 205, 124, 306, 228, 1210 (underlined for 1] 10), and 239. The strongest peaks for the phase were 237 and 330 but they were strongly influenced by the strong peak of abcc 532, 611, and the very strong peaks of a-Al 200 and a-Fe 110. aH 236 (d= 214.1 pm), 501 (d = 213.6 pm), and 3010 (d =211.0 pm) as listed in the JCPDS file could not be used as they
are strongly influenced by abcc 433, 530, and h 002. Not all the diffraction peaks listed in the JCPDS file for aH phase could be detected in pattern OC1 (Fig. 3). Each JCPDS data file is based on a diffraction experiment using a sufficient amount of powder. In the present case the outer compact layer is very thin, as shown in Fig. 1. Hence, it was reasonable for many weak diffraction peaks as listed in JCPDS to be absent. On the other hand, JCPDS for the aH phase only lists diffraction peaks with d-spacings \ 170 pm (2u = 54°). This is why the weak but distinctive peak at 2u= 73.54° in pattern OC1 has been labeled as UI (unidentified). For aH 3511 (d=128.7 pm and 2u= 73.52°) and aH 457 (d= 128.8 pm and 2u= 73.53°), the multiplicity factor is high (=24). For a hexagonal structure, the structure factor for 3511 and 457 is also high ( = 0.75f 2 where f is the scattering factor). Hence, it was likely that the peak at 2u= 73.54° was due to aH 3511, 457. In pattern OC2 (Fig. 4 and Table 4), diffraction peaks from the abcc phase, silicon and a-Al were absent meaning that the surface of the sample contained very little composite layer materials. This also suggests, as already discussed in Section 2.2, that only a very small amount of the outer compact layer was left on the sample surface and hence the intensities from aH diffraction were expected to be weak. Peaks from aH 112, 204, 205, 124, 137, 239 could be detected in the pattern. Due to the absence of abcc 532, 611 in pattern OC2, aH 237, which is one of the strongest diffraction peaks for the phase, was clearly identified. Furthermore, the aH 330 peak could also be identified.
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193
Table 4 Observed and calculated a (a= 1237.2 pm, c = 2615.8 pm) d-spacings, X-ray diffraction intensities for the phase isomorphous with aH-Fe2SiAl8 in patterns OC1 and OC2 and from JCPDS dobs (pm)b
dcal a(pm)
IOC1c
IOC2c
IJCPDS
558.9
559.2 535.7 524.7 495.8 456.5 449.4 436.0 414.4 400.2 386.9 374.3 367.3 357.2 352.8 344.3 330.5 337.0 320.2 306.5 296.7 295.0 289.8 289.1 279.6 276.3 274.6 270.5 267.9 255.5 254.4 252.3 245.6 244.7 241.2 238.4 236.6 233.8 232.6 228.2
vw
vw
14 8 16 7 10 2 8 12 44 5 12 6 16 19 34 16 2 28 1 4 5 2 2 3 5 1 2 2 2 7 3 8 3 2 4 10 36 18 9
414.3 399.9 374.3
344.4
320.2 296.5
276.2
245.8
232.4
w —d
vw vw
w
vw
m
mw
—e
—e
—f
vw
—g
mh
w
hkl 112 200 201 202 203 114 006 204 121 122 205 123 300 107 124 303 008 125 207 216 305 132 118 224 306 127 134 400 209 128 226 136 231 308 405 233 140 137 406
dobs (pm)
224.9 222.3 219.7
206.4 205.3
199.5
187.5
185.9
177.6
dcala (pm) 225.4 224.7 222.4 219.7 218.0 217.4 214.1 213.6 211.6 211.0 208.1 207.2 206.2 205.4 203.6 201.9 199.7 198.3 197.9 197.2 192.3 190.4 188.4 187.7 186.8 186.4 186.1 185.7 184.6 184.1 180.6 178.6 177.8 176.9 176.2 171.6 170.1 169.7
IOC1c
IOC2c
w —i w
—j —k —k —k
—k
—l sm —n
mwl mm —n
mwo
m
w
mwp
vw
w
IJCPDS
hkl
4 4 13 6 26 39 44 31 7 80 4 11 100 98 49 3 12 30 27 7 5 9 13 28 4 8 12 13 3 2 2 2 7 2 1 1 2 2
309 228 235 1210 0012 2011 236 501 502 3010 503 408 330 237 504 2012 2210 505 3011 243 506 152 2013 239 0014 336 3012 1311 154 1014 155 600 4011 602 430 520 434 2114
a
Based on the observed d-spacings as bolded. d-spacings bolded identify the phase with high confidence. c Intensity: vs, very strong; s, strong; m, medium; mw, medium to weak; w, weak; vw, very weak. d Overlaps with medium-strong abcc 310 peak. e Overlaps with medium-strong h 111 peak. f Overlaps with weak abcc 411 330 peak. g Overlaps with weak abcc 431 510. h Overlaps with weak Al 111. i Overlaps with weak abcc 440 peak. j Overlaps with medium-strong abcc 433 530 peak. k Overlaps with very strong h 002 peak. l Overlaps with strong aH 237 peak and very strong peak of combined Al 200 and Fe 110 in OC1 but more distinctively identified in OC2. m Overlap with very strong Al 220 and Fe 110 peak in OC1 but not in OC2. n Overlaps with strong abcc 532 611 peak in OC1 and strong Fe 110 in OC2. o Overlaps with weak abcc 620 peak. p Overlaps with weak abcc 631 peak. b
Richards et al. [13] recently used X-ray diffraction to identify the structures of intermetallic phases present in aluminized coatings. They identified, for hot dipping of
plain carbon steels in an Al–11Si melt, the outer intermetallic layer overlying the h-Fe2Al5(Si) layer to be aH-Fe2SiAl8. This is consistent with our identification of
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Table 5 Observed and calculated (a= 766.2 pm, b = 637.5 pm, and c =425.7 pm) d-spacings, X-ray diffraction intensities for the phase isomorphous with h–Fe2Al5 present in layers patterns OC2 and IC-H13 dobs. (pm)
dcal. (pm)
hkl
IOC2
IIC-H13
321.4 213.0 195.2 177.0 147.6 136.1 122.7 121.8
321.4 212.9 195.2 177.0 147.6 136.3 122.8 121.8
111 002 112 022 132 113 223 313
m vs m w vw w mw mw
mw s mw – – – – –
the outer compact layer to be aH phase, although detailed comparison could not be made, as Richards et al. did not give their detailed diffraction patterns. The abcc phase has not been reported to be present in aluminized coatings [13]. aH-Fe2SiAl8 is a stable Fe– Al–Si ternary phase [14] and the absence of the abcc phase is to be expected when aluminizing of plain carbon steels in an Al – Si – (Fe) melt.
3.3. Structure of the inner compact layer Strong diffraction from the inner compact layer was to be expected in patterns OC1 and OC2, as already explained in Section 2.2. As discussed, diffraction from abcc, silicon, a-Al of the residual composite layer and aH from the outer compact layer are found in pattern OC1 while only diffraction peaks of medium to weak
intensity from aH phase are present in pattern OC2. Hence, identification of the inner compact layer using pattern OC2 (Fig. 4) was simpler and clearer. Table 5 lists the data for pattern OC2 excluding those diffraction peaks due to the aH phase and the H13 substrate. From Fig. 4 and Table 5, eight diffraction peaks were identified for the structure isomorphous with orthorhombic h-Fe2Al5. This structural identification of the inner compact layer is consistent with that of the layer formed next to the steel substrate during aluminizing using a Al–11Si melt as a phase structurally isomorphous with h-Fe2Al5 [13]. Diffraction from the h-phase in patterns OC1 (Fig. 3) and OC2 (Fig. 4) are similar. All the h-Fe2Al5 diffraction peaks listed in Table 5 can be identified in pattern OC1. Further positive identification of the inner compact layer being a phase isomorphous with the h-Fe2Al5 was provided by the IC-H13 sample (Fig. 5 and Table 5). This diffraction pattern shows 111, 002, and 112 from the h-phase which was the residual material (of the inner compact layer) on the H13 immersed sample. From pattern IC-H13, where no diffraction peaks from the aH phase are present, the intensity of 111 is comparable to that of 112. For this reason, it can be concluded that in pattern OC2 (Fig. 4), the peak at 2u= 27.81° whose intensity is also comparable to h 112 in that pattern, was mainly contributed by h 111. Figs. 3–5 and Table 5 show that 002 from the h-phase had a very strong intensity, indicating that the phase has a preferred growth direction such that (001) of the phase was parallel to the steel substrate. This is
Fig. 5. X-ray diffraction pattern (scanning started from 20°) of the H13 sample (IC-H13 in Fig. 1) showing some amount of h-Fe2Al5 phase present on the sample surface.
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Fig. 6. EDS analysis of the intermetallic phases plotted as a function of the distance from the h – aH interface. Selected elements from the analyses on H13 are also plotted.
in common with the preferred orientation of the hFe2Al5 phase found in hot dipping of steel in aluminium melts or melts containing aluminium. The phase grows in the 001 direction towards the steel substrate in a variety of dipping conditions [13,15,16].
3.4. Compositional features of the intermetallic phases Fig. 6 shows a summarized EDS analysis on the intermetallic phases. For the h-phase layer, the average composition for major elements was 51% Al, 42% Fe, 4.5% Si, and 2.5% Cr. This compares to :52%Al, 45%Fe, and 3%Si for the h-phase layer formed during hot dipping of mild steel in an Al – 8%Si melt as recently given by El-Mahallawy et al. [17]. It seems that, in the present case, chromium may replace some iron in the h-phase. In El-Mahallawy et al. samples, precipitates rich in silicon (\ 10%) were clearly present in their h-phase layer. It is not known if these precipitates formed during hot dipping or during the slow cooling (in air) of the samples. Precipitation of the silicon-rich phase in the h-phase layer may have resulted in a decrease in silicon content in the h-phase in Mahallawy et al. samples. From the analytical data presented in Fig. 6, it appears that the chromium content (: 1.5%) was generally lower in the thin aH phase layer than that of the inner h-phase layer. It reverses to a higher content at
approximately 3% in the abcc phase in the composite layer and the iron content decreased slightly from aH to abcc phase. Manganese and copper contents were also lower in the aH phase than in the abcc phase. These compositional changes are consistent with the established phenomenon that chromium, manganese, and copper are the elements required to stabilise a cubic structure (abcc phase) at the expense of the hexagonal structure (aH-Fe2SiAl8) [9,12]. 4. Conclusion The three reaction layers formed during immersion of H13 tool steel in a die casting alloy melt have been structurally identified. The irregular intermetallic phase in the thick composite layer away from the H13 steel substrate was identified to have a bcc structure, abcc(FeSiAlCrMnCu). The thin and continuous layer between the composite layer and the inner compact layer was found to be structurally isomorphous with aHFe2SiAl8. The compositional difference between aH and abcc phases was mainly that the latter consisted of a higher amount of Cr+ Mn+Cu and a lower amount of iron. It is likely that the presence of chromium, manganese, and copper in the system resulted in the transformation aH abcc. The inner compact layer next to the steel substrate was identified to be orthorhombic h-Fe2Al5 containing silicon and chromium.
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Acknowledgements We would like to thank the Co-operative Research Centre for Alloy and Solidification Technology and Nissan Casting Australia for supporting this work. We would also like to thank N.M. Rockelmann and Dr D.G. Hay for the production of X-ray diffraction patterns.
[6] [7] [8] [9] [10]
[11]
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