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Erosion and corrosion of alumina refractory by ingot casting steels
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Erosion and corrosion of alumina refractory by ingot casting steels Jens Fruhstorfer a,∗ , Leandro Schöttler b , Steffen Dudczig a , Gert Schmidt a , Patrick Gehre a , Christos G. Aneziris a a b
Institute of Ceramic, Glass and Construction Materials, TU Bergakademie Freiberg, Agricolastraße 17, 09596 Freiberg, Germany Deutsche Edelstahlwerke GmbH, Obere Kaiserstraße, 57078 Siegen, Germany
a r t i c l e
i n f o
Article history: Received 18 August 2015 Received in revised form 12 November 2015 Accepted 22 November 2015 Available online xxx Keywords: Contact angle Manganese Aluminum Silicon Oxygen
a b s t r a c t This study investigated the corrosion of alumina-based refractories by alloyed steels. The alumina crucibles were physically characterized. Four steels with varying casting temperature and alloy contents were melted under argon atmosphere and the contact angles measured. Post-mortem, X-ray diffraction, optical and scanning electron microscopy and energy-dispersive X-ray spectrometry were used to analyze the erosion and corrosion. The refractory properties did not affect erosion and corrosion. No infiltration occurred. The contents of oxygen and deoxidizing steel alloys influenced the corrosion and erosion of the matrix. Firstly, manganese reacted to manganese aluminate. Silicon followed—forming manganese alumosilicates. For highest casting temperatures (1580 ◦ C) also mullite formed. At the steel–refractoryinterface aluminum displaced the silicates, enhancing corrosion. Due to mismatching thermal expansions, spalling occurred for the steel 17CrNiMo76 respectively 1.6587 in which mullite formed. In the coarse fraction only chromium doped alumina formed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction A new generation of secondary steel metallurgy developed [1] as well as the demand for improved steel qualities is continuing. The quality requirements are met by improving the steel cleanliness by reducing non-metallic inclusions. Non-metallic inclusions are a significant problem because they act as stress raisers in the cast steels and, thus, limit the attainable mechanical properties [2,3]. During the commonly applied continuous casting of steel severe bending stresses occur at high temperatures. Steel grades which are sensitive to such stresses as well as steels for special applications are ingot cast discontinuously. These steels may be alloyed with carbon, chromium or manganese [3–7]. Ingot casting is usually performed in a bottom teeming process by pouring the steel through a refractory pipe system into the ingot mould [3,5]. According to Zhang and Thomas [2,3] non-metallic inclusions in steel are classified into indigenous and exogenous inclusions. Indigenous inclusions are deoxidation products or inclusions which precipitate during cooling. Deoxidation products can be e.g. Al2 O3 in aluminum killed steels, SiO2 in silicon killed steels or MnO in steels containing manganese [8]. Exogenous inclusions are
∗ Corresponding author. E-mail address: [email protected] (J. Fruhstorfer).
formed by chemical and mechanical interaction of the steel with its surroundings. They have, generally, the most harmful effects on the mechanical properties due to their larger defect size and surface-near location. Sources of exogenous inclusions are mainly reoxidation, slag entrainment, lining erosion and chemical reactions. Regarding the boundary interaction of an ingot cast steel with its refractory pipe system, the most serious inclusions arise from lining erosion and chemical reactions. Refractory contents with a low oxidation potential are reduced by steel elements with a high oxidation potential as e.g. Al can reduce SiO2 leading to Si and Al2 O3 [8]. Therefore, refractories containing SiO2 provide a high oxidation potential. Besides the reoxidation, compound formation is a source of additional non-metallic inclusions. In Al-killed steels containing Mn, manganese aluminates [9] and in Si-killed steels containing Mn, manganese silicates can be forming. In the complexer MnO–SiO2 –Al2 O3 system, manganese alumosilicates may be forming [2,3,10]. Also Cr-alloyed steels are reported to be highly corrosive [5–7]. Chromium oxide is forming solid solutions with alumina [11]. Furthermore, it is able to react to chromite [12,13] or alumino-chromite spinels [13]. Despite the manifold possible reactions, it was reported that high purity alumina, mullite or zirconia parts with a low amount of free silica showed a good applicability and resistance to Mnand Al-rich steels. Furthermore, they have an excellent hot erosion resistance [2,3,5].
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.11.038 0955-2219/© 2015 Elsevier Ltd. All rights reserved.
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2 Table 1 Steel compositions and casting temperatures. Steel grade
1.3520 1.3816 1.4542 1.6587
Casting temp.
1522 ◦ C 1460 ◦ C 1535 ◦ C 1580 ◦ C
Table 2 Large crucible firing curve.
Chemical composition in wt% C
Si
Mn
Cr
Al
0.93 0.07 0.03 0.16
0.60 0.35 0.40 0.21
1.10 19.05 0.70 0.63
1.50 18.17 15.30 1.52
0.00 0.01 0.01 0.03
However, the corrosion and wear experiments were conducted with steels varying in their composition. Although some reactions of alloys with contents of the refractory were determined, the interaction of the alloys regarding their combined effect on the corrosion is rather unclear. Consequently, no highly corrosive steel for corrosion tests for comparing different refractory materials was determined. The purpose of this study, therefore, is to investigate the corrosion and erosion of cast alumina crucibles by four ingot casting steels of varying composition and casting temperature. A highly corrosive steel for further and future corrosion tests shall be determined. In this study, the corrosion of the alumina crucibles will be regarded after use as an indirect potential of producing inclusions and to evaluate the corrosion mechanisms [14]. Erosion and corrosion will be analyzed microscopically, chemically and mineralogically. 2. Experimental Corrosion tests with four steels varying in composition and cast alumina crucibles were carried out in an inductively heated metal casting simulator. The tested ingot casting steels were a C- and Si-rich (1.3520 resp. 100CrMn6), a Mn- and Cr-rich (1.3816 resp. X10MnCrN1818), a Crrich (1.4542 resp. X5CrNiCuNb174) steel and an Al-rich steel with a high casting temperature (1.6587 resp. 17CrNiMo76). Their chemical compositions as measured by X-ray fluorescence spectroscopy (ARL 9800, Thermo Scientific, U.S.A.) and combustion analysis (CS 244, LECO Corporation, U.S.A.) for the determination of the carbon content as well as their casting temperature are presented in Table 1. The raw materials for the refractory crucibles were alumina fractions up to 3 mm (Tabular Alumina T60/64), reactive alumina (CL370) and a re-hydratable alumina binder (Alphabond 300) obtained from Almatis GmbH, Germany. Additionally, a defoamer (Contraspum K 1012, Zschimmer & Schwarz GmbH & Co. KG, Germany), a dispersant (Castament FS 60, BASF Construction Polymers GmbH, Germany) and water were used, described in more detail by Fruhstorfer and Aneziris [15]. The composition of the castable for manufacturing large crucibles with a height of ≈25 cm and a diamter of ≈27 cm was based on a modified Andreasen model as introduced by Fruhstorfer and Aneziris [15]. In the modified model, the particle size distribution was described by a minimum and a maximum distribution modulus. In this study a maximum distribution modulus of 0.8 similar to the previous study [15] was used. In difference, a minimum modulus of 0.27 and a water content of 5.5 wt% were applied because preliminary experiments showed that with minimum moduli ≤0.27, maximum ones ≥0.8 and slightly increased water contents the flowability increased by simultaneously further decreasing the pore sizes arising during shaping. The binder content was adjusted to 3 wt%. Furthermore, 0.1 wt% of the defoamer and 0.57 wt% of the dispersant related to the solids content were added. The castables of about 21 kg mass were prepared with a compulsory mixer (R05T, Eirich, Germany). The batches were dry mixed
ϑstart in ◦ C
ϑend in ◦ C
Time in h
ϑroom 180 600 800 1600 1600 600
180 600 800 1600 1600 600 ϑroom
3 42 4 7 3 13 Free cooling
for 1 min, then the water was added and it was wet mixed for 5 min. Right after finishing mixing, the crucible moulds were filled bottom up. The moulds had a silicone covered core and were lubricated with a Teflon oil. The open top of the mould was covered with a wet tissue after complete filling. After setting for 1 d, the crucible was demoulded. The large crucibles were dried carefully due to the strong dehydration of the re-hydratable alumina binder and fired as presented in Table 2. After firing the filling volume of the crucibles was about 7.5 L. The mineralogical composition was determined by Xray diffraction (XRD) (PW1820, Philips, Netherlands) with Cu-K-␣ radiation. The corrosion experiments were conducted in an inductively heated melting unit (150 kW inductor power) of a metal casting simulator (Systec GmbH, Germany) introduced by Dudczig et al. [16,17] and Aneziris et al. [18]. The crucibles were filled with 20–25 kg steel. To reduce the oxygen content it was twice evacuated to 2.5 mbar. After the first evacuation it was flushed with argon 4.6 gas (99.996 vol% argon) to a pressure of 450 mbar and after the second evacuation to ambient pressure. Then it was inductively heated by increasing the power by 10 %/h related to the 150 kW. This resulted in an interpolated heating rate of ≈250 K/h. The steels were heated to their casting temperature (Table 1) and hold at this temperature for 1 h because in steel ingot casting the contact time before solidification is about 15–45 min in line with Zhang and Thomas [3]. By applying an hour contact time in an induction heating process the corrosion effect was maximized. Every 10 min, the temperature and oxygen content were measured with a pO2 /T-sensor-system. Cooling was carefully performed under argon atmosphere by reducing the power by 5 %/h to reduce the risk of cracking. The benefit from the large scale experiments was the resembling of the application and the obtainment of enough corroded material for detailled analysis. The crucibles were cut and analyzed after the corrosion test. The spots where mainly parallel steel flow occurred were analyzed because in the refractory pipe system also mainly parallel flow occurs. The thickness of the eroded layers was analyzed by digital optical microscopy (DOM) (VHX Digital Microscope 2000, Keyence Microscope Europe, Belgium) and the corroded layer visually or by microscopy. In the same spots on the outer area of the crucible as well as on spots of the upper crucible part, porosity and pore sizes below 150 m were measured as described in a previous study [15] with the water immersion method and mercury porosimetry, respectively. The mercury porosimetry was analyzed by the Washburn method applying an assumed contact angle of 140◦ and a mercury surface tension of 0.48 N/m. From the reacted areas, also samples for scanning electron microscopy (SEM) (ESEM XL30FEG, FEI company, Netherlands) equipped with an energy-dispersive X-ray spectrometer (EDX) (EDAX, Ametek GmbH, Germany) were prepared and analyzed. The SEM signal of back-scattered electrons was used to visualize phase differences due to their differing Z-contrasts [19,20]. To determine the mineralogy, XRD was performed. Therefore, the visually corroded sections respectively layers of the crucible were separated from each other by crushing them into pieces and classifying them according to
Please cite this article in press as: J. Fruhstorfer, et al., Erosion and corrosion of alumina refractory by ingot casting steels, J Eur Ceram Soc (2015), http://dx.doi.org/10.1016/j.jeurceramsoc.2015.11.038
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3
Table 6 Pearson correlation coefficients of the mean oxygen content.
Total porosity
Open porosity
Bottom
Top
Bottom
Top
1 (1.3520) 2 (1.3816) 3 (1.4542) 4 (1.6587)
19.9 20.7 20.9 19.9
19.4 20.5 20.7 19.6
16.1 17.1 17.5 16.2
15.7 16.9 17.5 15.8
Average Total
20.4
20.1
16.7
16.5 16.6 ± 0.8
20.2 ± 0.6
Table 4 p-Values of the ANOVA for the total and open porosity (significant effects on a level ≤0.05 marked bold). Factor
Crucible number Location
p-Value of the Total porosity
Open porosity
0.0015 0.0142
0.0029 0.0955
Significant Bonferroni-corrected correlationon a level ≤ 0.01. on a level ≤ 0.05. on a level ≤ 0.1.
their apparent color. After classification, the pieces were milled to particle sizes below 63 m for the XRD measurement. Furthermore, the contact angles between cut slides of sintered parts and the different steels were measured by the sessile-drop method at the casting temperatures. A hot stage microscope as described by Aneziris and Hampel [21] was used. The casting temperature was hold for 1 h and it was continuously flushed with argon gas.
on a level > 0.1.
practical considerations in line with Schafföner and Aneziris [24] where comparable results were obtained for pressure slip castables. The pore size describing parameters were d50 = (2071 ± 111) nm and d80 = (5338 ± 1675) nm in the pore size range <150 m. The di means that i % of the pores were smaller than this size. The samples were taken from the bottom location. In dependence on the crucible resp. the porosities, no significant differences of the pore sizes were detected. In comparison to the data of Fruhstorfer and Aneziris [15] (d80 = 3691 nm) the d80 pore size tended to increase slightly. It was found in the previous study that for denser castables the pore sizes shifted towards smaller values, increasing their di in the pore size range <150 m. Consequently, the overall pore sizes were further shifted towards smaller values by using slightly more fines and 0.5 wt% more water.
3. Results and discussion Firstly, the refractories and, afterwards, the melting tests were investigated. Finally, the corrosion results were correlated with the refractory and steel properties. 3.1. Alumina refractory characterization
3.2. Melting tests
The qualitative XRD analysis of the crucibles detected only corundum and traces of -alumina after firing. The total and open porosities of the crucibles are listed in Table 3. The total porosity was (20.20 ± 0.59) vol%. In a previous study [15], 5 wt% water were added to a self-flowing castable resulting in an average total porosity of 19.4 vol% and 15.7 vol% open porosity. By using 5.5 wt% water in the current investigation, each porosity increased for about 1 vol%. To investigate whether the hydrostatic pressure causes a density gradient, an analysis of variances (ANOVA) was used to determine the main effects of the factors ‘Crucible number’ and ‘Location’ (bottom or top) on the total and open porosity. An ANOVA tests means on their statistical difference [22]. The tests were performed with the statistics software R [23]. An effect was regarded significant when its p-value was ≤0.05. Table 4 presents the resulting p-values. As the factor ‘Location’ had a significant influence on the total porosity, in the cast material a density gradient arose. However, the influence of the ‘Crucible number’ was more significant. Thus, the statistical differences of the factor ‘Location’ were negligible for
During the holding time, every 10 min the temperature and oxygen content were logged. The averages and standard deviations are presented in Table 5. The specified casting temperatures were met satisfactorily. The highest inaccuracy showed the melting test of the Mn- and Cr-rich steel 1.3816—the average temperature was about 10 K below the required one. The arising oxygen content was correlated with the average melting temperature as well as with the steel and crucible properties. Table 6 presents the Pearson product–moment correlation coefficients [25] and the p-values of the t-tests on difference from zero [22,23]. In the present study, coefficients below 0.5 are regarded to indicate a poor linearity, coefficients between 0.5 and 0.8 a medium one and larger coefficients a strong linear correlation. It can be seen in Table 6 that the crucible properties (porosities and pore sizes) showed no significant correlations with the oxygen content. The determined significant differences between the crucibles (see Table 4 in the prior subsection) had no practical influence
Table 5 Logged data during steel melting in the large crucibles (* specified by steel manufacturer). Steel grade
Casting temp.*
Logged temperature in ◦ C
1.3520 1.3816 1.4542 1.6587
1522 ◦ C 1460 ◦ C 1535 ◦ C 1580 ◦ C
1519 1451 1536 1580
± ± ± ±
9 26 6 10
Logged oxygen content in ppm 10 1 20 36
± ± ± ±
1 2 1 2
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Fig. 1. Dependency of the oxygen content in the melt on the highly significant correlation properties.
on the oxygen content. But the temperature showed a significant correlation with a strong correlation coefficient indicating a high linearity of the relation (see Fig. 1(a)) similar to the observations of Kim and McLean [9]. Manganese, silicon and chromium are possible deoxidants [8] and, therefore, their correlation coefficients were negative and significant. The correlations were medium because the relations were supposedly not linear due to an exponential nature or occurring interactions—in the present study the relation on the manganese content was exponential as can be seen in Fig. 1(b) similar to the results of Kim and McLean [9]. The dependency of the oxygen content in the melt on the aluminum content was positive and significant although Al is a known strong deoxidant. Therefore, the observed relation is presumably indirectly determined. In Fig. 1(c) can be seen that at a level of 0.01 wt% Al two levels of resulting oxygen contents (≈3 ppm and ≈20 ppm) arose. The lower oxygen content resulted when the Mnand Cr-rich steel 1.3816 was melted at ≈1451 ◦ C—Al acted deoxidizing. The higher oxygen content arose during melting of the Cr-rich steel 1.4542 at ≈1536 ◦ C. The aluminum amount interacted with the temperature. For higher temperatures Al presumably affected the corrosion significantly and, therefore, did not show a deoxidizing behavior. 3.3. Corrosion analysis Wetting, erosion and corrosion were analyzed. Firstly, the wetting was analyzed by regarding the contact angles. An ANOVA showed statistically that the contact angle did
not change during 1 h holding time. Therefore, the angles at 0 and 60 min were used for the analysis. The values are presented in Table 7. Contact angles above 90◦ characterize a non-wetting behavior and contact angles above 110◦ a non-reactive one according to Eustathopoulos et al. [26]. It was, therefore, tested with one sample Bonferroni-corrected t-tests if the measured contact angles were greater than 90◦ and lower than 110◦ . Only the Mn- and Cr-rich steel 1.3816 wetted the alumina ( < 90◦ ) and only the C- and Si-rich steel 1.3520 showed a non-reactive behavior ( ≥ 110◦ ). The correlations of the contact angle with the steel, crucible and melting properties are presented in Table 8. On a significance level of 0.05 after Bonferroni correction, only the manganese content and the d80 pore size correlated significantly negative—leading to decreased contact angles. In line, an increasing manganese content of iron–manganese alloys on solid alumina under argon atmosphere at 1560 ◦ C led to decreasing contact angles [10]. An increasing d80 means that the sample contained less small pores which contribute to the surface area. And a decreasing surface area commonly contributes to an increasing advancing contact angle [26]. The corrosion was firstly analyzed macroscopic, regarding the depth of the eroded and corroded section including the eroded part. Tables 7 and 8 present the values and the correlation coefficients. The eroded section correlated on an adjusted significance level of 0.05 with no measured property. Highly probable interactions between the properties on the erosion could cause this non-significance. Therefore, also the coefficients significant on a
Table 7 Large crucible corrosion test results.
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Table 8 Pearson correlation coefficients.
Fig. 2. SEM image of alumina corroded by C- and Si-rich steel 1.3520.
Significant Bonferroni-corrected correlation. on a level ≤ 0.01. on a level ≤ 0.05. on a level ≤ 0.1. on a level > 0.1.
level of 0.1 were marked in Table 8. The erosion, thus, correlated with the oxgen and aluminum content. Both correlations showed a medium linearity affirming an interactive behavior. An increasing oxygen content of the melt led to increased erosion in line to the reviewed results by Zhang and Thomas [3] where the fatigue lifetime of bearing steels decreased exponentially with increasing oxygen contents. For every corrosive attack of liquid steel, oxygen was a reactant—independent of the other reactants. Furthermore, increasing amounts of Al thickened the eroded layer. Supposedly, after the first reactions between the crucibles and the other elements, the Al reacted with the first corrosion product—decomposing it or forming a more complex product. Both iterated corrosion mechanisms would explain the non-deoxidizing behavior observed in the prior subsection. The complete corrosion correlated with the silicon content negatively on a significance level of 0.1 in addition to the oxygen and aluminum content. As silicon acted deoxidizing (see prior subsection), it reduced the oxygen amounts and, consequently, the corrosion. But the other deoxidants like Mn did not show significant influences, although their oxidation potential is higher than the one of silicon [8]. Possibly, the other deoxidants reacted with the refractory in elementary or oxidic form. Oxidized silicon has also a high potential to form crystalline or amorphous silicate structures with manganese oxide or alumina. Anyway, such subsequent reactions would not necessarily increase the extent of the complete corrosion. Regarding the aluminum effect, the presumption from the prior subsection that it strongly affects the corrosion was confirmed. The XRD results corresponded to the colors of the layers. In the crucible with no visible layer, corroded by the C- and Si-rich steel 1.3520 only alumina was detected by XRD. The non-reactive behavior confirmed the wetting result. The corroded layer of the Mn- and Cr-rich steel 1.3816 had a tinted yellow color, corresponding to manganese aluminate. Additionally, some amorphous parts existed. The amorphous parts were qualitatively evaluated by regarding the peak heights of the chromium doped alumina in combination with the following presented EDX measurements. In case of the Cr-rich steel 1.4542, the corroded layer was parted into an inner (near steel) red–violet layer of about 1 mm thickness and
a subsequent tinted yellow layer. The red–violet layer contained mainly chromium doped alumina and traces of compounds of the system Mn–Si–Cr–O. It presented also the largest amorphous parts. Of the tinted yellow layer not enough material for a separate XRD analysis was obtained. Anyway, as for the complete corroded section also manganese aluminate traces were found which were not detected in the red–violet layer, the tinted yellow layer contained manganese aluminate. The eroded respectively partly spalled layer of the Al-rich steel 1.6587 with the highest casting temperature was tinted red colored, corresponding to chromium doped alumina and few manganese aluminate with some amorphous parts. The subsequent corroded layer was again tinted yellow due to manganese aluminate but had in contrast to the Mn- and Cr-rich steel 1.3816 very low amorphous parts. The XRD study of the individual layers returned also that mullite formed in the tinted yellow section of the alumina corroded by 1.6587. Between the mullite containing tinted yellow layer and the prior tinted red one containing chromium doped alumina, a thermal expansion mismatch existed. This mismatch led to spalling respectively delamination and explained the strongly increased erosion (see Table 7). This is in line with a study reviewed by McCauley [27] of a SiC–Al2 O3 composite, where upon oxidation at 1530 ◦ C a zone containing mullite and amorphous alumo-silicates formed. Delamination was reported as a possible consequence. To determine the corrosion reactions in more detail, SEM and EDX measurements were analyzed. The results of the EDX measurements are presented in Table 9. The SEM image in Fig. 2 presents a section of the alumina crucible corroded by the C- and Si-rich steel 1.3520. Although by XRD only alumina was detected, in the image additional phases were observed. The bright areas contained mainly iron oxide, which was found to a depth of about 40m. The gray areas corresponded to manganese aluminate found to a depth of about 30m. Down to a depth of about about 90m rarely iron aluminates were found. The porosity was not filled with steel—infiltration did not occur. The SEM image in Fig. 3(a) presents a section of the alumina refractory corroded by the Mn- and Cr-rich steel 1.3816. The XRD results showed, that only manganese aluminate was detected. The EDX measurements presented in Table 9 were in line and visualized for the matrix (finest particles) in Fig. 3(b) dependent on the depth. A measurement on a coarse grain showed that it reacted only to a small extent. Only to a depth of about 300m, additionally, chromium resp. chromium oxide doping was detected in amounts below 0.5 mol%. With increasing depth, the content of manganese oxide of the manganese aluminate compositions decreased—the non-stoichiometry explains the increased amorphous parts observed by XRD. Additionally, it can be said that the corrosion rate was controlled by diffusion. The aluminate compositions were detected to a depth of 1.4 mm. Silicon contents in the aluminate were detected to a depth of 1 mm. Supposedly, it
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followed after the manganese reacted with the refractory. It appeared in the non-stoichiometric and presumably amorphous aluminate structures and hence was not detected by XRD. This affirmed the presumption of the iterated corrosion. The refractory was not infiltrated by the steel. In Fig. 4 the SEM image of an upper section of an alumina sample corroded by the Cr-rich steel 1.4542 is presented as well as EDX measurements dependent on the depth. At the interface to the steel, only chromium doped alumina was found to a depth of about 100 m. Subsequently, a layer of manganese alumo-silicates followed to a depth of 200m while larger particles corresponded to chromium doped alumina. In the following section, chromium doped manganese aluminates were found besides chromium doped manganese alumosilicates. The chromium contents decreased with increasing depth till 800m. For depths below ≈2.1 mm only manganese aluminates were found—from there also the tinted yellow layer started. The red–violet layer formed due to the corrosion. Although, many different compositions were detected by XRD, from the EDX images is clear that also many amorphous non-stoichiometric structures formed. No infiltration was observed. Firstly, it seems, the manganese reacted with the alumina refractory. The increased casting temperature led to an increased crystallization of the compound compared to the manganese aluminate formed during corrosion by the Mn- and Cr-rich steel 1.3816. Subsequently, the silicon reacted with the refractory and the manganese aluminate, forming silicates besides the manganese aluminate. Towards the interface to the steel, the aluminate was gradually displaced by silicon containing reaction products. Anyway, at the interface to the steel only chromium doped alumina was detected. Consequently, the silicates were displaced at the interface by the aluminum at higher temperatures. Thus, the aluminum decomposed the previously formed corrosion products. The freed educts like manganese oxide and silica were hence available for new corrosion reactions. The aluminum enhanced the corrosion
Fig. 3. Alumina corroded by Mn- and Cr-rich steel 1.3816.
Table 9 Average EDX results for different depths and probable corresponding compounds Steel grade
1.3520
1.3816
1.4542
1.6587
Affected particles
Depth in m
Finest Fine Finest Finest Coarse Finest Finest Finest Finest Fine Finest Finest Finest Fine Finest Finest Finest Finest Finest Finest Finest Coarse Fine Finest Finest Finest Finest Finest Finest
20 20 80 300 300 700 1000 1400 70 150 160 310 310 550 560 800 1300 2050 2050 3200 50 50 350 525 700 1400 2500 5000 6300
Chemical composition in mol% Al2 O3
MnO
MgO
SiO2
Fe2 O3
Na2 O
Cr2 O3
Mo2 O3
CaO
52.2 2.1 61.5 49.9 96.6 71.1 74.2 95.7 93.5 94.8 43.0 33.2 54.0 94.9 46.6 42.3 38.8 56.9 40.2 59.2 34.7 98.7 48.0 17.5 17.9 53.3 52.3 60.0 52.1
41.0 0 0 48.6 1.3 17.3 10.1 1.5 0 0 28.1 44.7 7.9 0 42.7 10.0 10.0 36.2 15.3 34.7 7.7 0 36.8 16.9 17.7 41.7 43.8 34.5 44.5
2.0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 3.9 0 3.1 0 0 10.3 0.9 0.9 1.8 1.4 0.4 0.8
0.9 0 0 1.1 2.1 5.8 8.3 1.6 0 0 14.1 0.5 33.1 0 0.9 40.6 43.4 0 37.1 0 46.6 0 0 51.2 50.9 0 0 2.1 0
2.9 91.8 24.2 0 0 0 0 0 0 0 0 0 0 0 1.5 0.3 0.3 3.1 0.6 3.0 1.0 0 0.9 0 0 2.3 2.3 1.7 2.6
0.5 3.2 7.8 0 0 5.7 6.8 1.2 0 0 0.8 0.7 3.3 0 0 6.4 7.6 0 4.4 0 4.9 0 0 7.4 7.7 0.5 0 1.4 0
0.3 0.3 0 0.4 0 0 0 0 6.5 5.2 14.0 20.9 1.7 5.1 8.4 0.5 0 0 0 0 0.2 1.3 3.7 0 0 0.6 0.2 0 0
0 2.5 1.2 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0
0 0 0 0 0 0.2 0.6 0 0 0 0 0 0 0 0 0 0 0 2.3 0 4.0 0 0 5.0 4.2 0 0 0 0
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Fig. 4. Alumina corroded by Cr-rich steel 1.4542.
therefore indirectly, which explained finally the correlations and confirmed the theory of the iterated corrosion. Fig. 5 shows the upper section of the alumina corroded by the Al-rich steel 1.6587 with the highest casting temperature and EDX measurements. The white ball-shaped structures were mainly iron (e.g. 99.1 at% Fe, 0.9 at% Ni). As no infiltration was observed, the iron balls were formed most likely by a reduction of iron compounds. In the SEM-Fig. 5(a) can be seen that the surface was eroded strongly for at least 200 m besides the observed spalling reported in Table 7. This is in the size range of the Cr-doped alumina layer
Fig. 5. Alumina corroded by Al-rich steel 1.6587 with the highest casting temperature.
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observed for the corrosion by the Cr-rich steel 1.4542. Presumably due to the erosion, the alumina layer was not found by EDX in Fig. 5(c). That it existed also here was supported by the increasing alumina contents for depths more near to the surface. The increased hot surface erosion might be caused by the high treatment temperature of 1580 ◦ C. In the coarse fraction at the surface-near locations Cr-doped alumina and in the fine but not finest fractions, manganese aluminate was found. While the chromium was solid soluted in alumina, this confirms the hypothesis of the iterated corrosion, starting with the manganese and following silicon because in the finest fraction to a depth of about 1 mm manganese alumino-silicate compositions were found. As they were not detected by XRD they were presumably amorphous. Only to a depth of 50m the silicates also contained chromium. The following layer was the tinted yellow one corresponding to mullite and manganese aluminate by XRD and chromium doped alumina in the coarse fraction. The 2:1-mullite layer formed in this sample between 1 and 1.5 mm depth because there the necessary silicon amounts were reached as visualized in Fig. 5(c). It was, therefore, situated behind the spalling region, confirming that in interaction with the temperature the mullite formation and the thermal expansion mismatch led to the spalling. Below a depth of 1.5 mm, mainly manganese aluminate spinels containing chromium were detected. Chromium was detected to a depth of 2.5 mm. The subsequent manganese aluminate spinel section reached to a depth of 6.3 mm. In contrast to the EDX results of the Mn- and Cr-rich steel 1.3816, no gradually decreasing manganese content was observed, explaining the higher crystallinity. Furthermore, it can be said that the corrosion rate by the steels 1.6587 and 1.4542 with higher casting temperatures was chemical reaction controlled. The chromium doping interacted with the temperature. For larger amounts and higher temperatures, the depth the chromium reached increased. Anyway, it seemed not to cause harmful damage to the refractory. The result that the Al-rich steel 1.6587 with the highest casting temperature led to the most harmful erosion and corrosion was conform to practical observations in Deutsche Edelstahlwerke GmbH steelworks. 4. Conclusion The erosion and corrosion of alumina crucibles by ingot casting steels of varying composition was investigated. • The contents of oxygen and the deoxidizing steel alloys significantly influenced erosion and corrosion. • An iterated corrosion occurred dependent on the temperature. The corrosion started with manganese reacting to manganese aluminate with the alumina refractory. Subsequently, a transition to manganese alumo-silicates occurred with silicon, firstly occurring besides the aluminates. At the interface, aluminum displaced these silicates in a sharp transition. Thus, the silica was again available for further corrosive reactions. Consequently, the aluminum content enhanced the corrosion strongly. • The highly corrosive steel 1.6587 (17CrNiMo76) led to the largest erosion and corrosion of the alumina-based refractories. Thus, it is recommended for further corrosion tests for comparing different refractory materials. Only during this corrosion test a mullite containing zone with a different thermal expansion formed due to the high casting temperature of 1580 ◦ C, leading to spalling. The high casting temperature and its interactions with the steel alloys, consequently, made this steel the most corrosive.
In a future study, other alumina- as well as mullite-based refractory materials will be tested in the steel casting simulator together with the recommended reference steel. Furthermore, also the interaction of the chromium with the temperature could be regarded in more detail as well as the effect of the porosity. Acknowledgment The authors thank the German Research Foundation (DFG) for supporting this collaborative study under grant numbers AN 322/27-1, AN 322/19-1 and AN 322/17-2. References [1] C.G. Aneziris, U. Klippel, W. Schärfl, V. Stein, Y. Li, Functional refractory material design for advanced thermal shock performance due to titania additions, Int. J. Appl. Ceram. Technol. 4 (6) (2007) 481–489. [2] L. Zhang, B.G. Thomas, Inclusions in continuous casting of steel., in: XXIV National Steelmaking Symposium, 2003, pp. 138–183. [3] L. Zhang, B.G. Thomas, State of the art in the control of inclusions during steel ingot casting, Metall. Mater. Trans. B 37B (2006) 733–761. [4] E. Fuchs, P. Jonsson, Inclusion characteristics in bearing steel before and during ingot casting, High Temp. Mater. Process. 19 (5) (2000) 333–344. [5] W. Ritter, K.G. Ruwier, W. Schönwelski, Carbonaceous fireproof material for use when casting steel in a bottom casting process and formed parts produced thereof, Patent WO 2011/054872 A1 (2011). [6] R.D.P. Rettore, E. Gueguen, W. Ritter, High alumina refractory material with improved thermal shock resistance and hot properties for hollowware, Refract. Eng. (November) (2012) 19–22. [7] W. Schönwelski, K. Ruwier, S. Foellbach, J. Sperber, Feuerfestkeramik für hochwertige Stähle beim Blockguss, Stahl Eisen 134 (9) (2014) 58–62. [8] H.J.T. Ellingham, Reducibility of oxides and sulphides in metallurgical processes, J. Soc. Chem. Ind. (1944) 125–133. [9] C.K. Kim, A. McLean, Thermodynamics of iron–manganese aluminate spinel inclusions in steel, Metall. Trans. B 10B (1979) 575–584. [10] Verein Deutscher Eisenhüttenleute (VDEh) (Ed.), Slag Atlas, 2nd ed., Verlag Stahleisen GmbH, Düsseldorf, 1995. [11] P.L. Roeder, F.P. Glasser, E.F. Osborn, The system Al2 O3 –Cr2 O3 –SiO2 , J. Am. Ceram. Soc. 51 (10) (1968) 585–594. [12] M. Tanahashi, N. Furuta, C. Yamauchi, T. Fujisawa, Phase equilibria of the MnO–SiO2 –CrOx system at 1873 K under controlled oxygen partial pressure, ISIJ Int. 41 (11) (2001) 1309–1315. [13] C.K. Kim, A. McLean, Thermodynamics of iron alumino-chromite spinel inclusions in steel, Metall. Trans. B 10B (1979) 585–594. [14] J. Poirier, F. Qafssaoui, J.P. Ildefonse, M.L. Bouchetou, Analysis and interpretation of refractory microstructures in studies of corrosion mechanisms by liquid oxides, J. Eur. Ceram. Soc. 28 (2008) 1557–1568. [15] J. Fruhstorfer, C.G. Aneziris, The influence of the coarse fraction on the porosity of refractory castables, J. Ceram. Sci. Technol. 5 (2) (2014) 155–166. [16] S. Dudczig, C.G. Aneziris, U. Ballaschk, Testing of thermal shock resistant fine grained alumina ceramics in a metal casting simulator, Refract. Worldforum 4 (1) (2012) 1–7. [17] S. Dudczig, C.G. Aneziris, M. Dopita, D. Rafaja, Application of oxide coatings for improved steel filtration with the aid of a metal casting simulator, Adv. Eng. Mater. 15 (12) (2013) 1177–1187. [18] C.G. Aneziris, S. Dudczig, J. Hubalkova, M. Emmel, G. Schmidt, Alumina coatings on carbon bonded alumina nozzles for active filtration of steel melts, Ceram. Int. 39 (2013) 2835–2843. [19] G.E. Lloayd, Atomic number and crystallographic contrast images with the SEM: a review of backscattered electron techniques, Mineral. Mag. 51 (1987) 3–19. [20] J. Fruhstorfer, S. Schafföner, C.G. Aneziris, Dry ball mixing and deagglomeration of alumina and zirconia composite fine powders using a bimodal ball size distribution, Ceram. Int. 40 (9) (2014) 15293–15302. [21] C.G. Aneziris, M. Hampel, Microstructured and electro-assisted hightemperature wettability of MgO in contact with a silicate slag-based on fayalite, Int. J. Appl. Ceram. Technol. 5 (5) (2008) 469–479. [22] D.C. Montgomery, Design and Analysis of Experiments, 5th ed., John Wiley & Sons, 2001. [23] R Development Core Team, R: A Language and Environment for Statistical Computing, R Foundation for Statistical Computing, Vienna, Austria, 2010. [24] S. Schafföner, C.G. Aneziris, Pressure slip casting of coarse grain oxide ceramics, Ceram. Int. 38 (2012) 417–422. [25] J. Benesty, J. Chen, Y. Huang, I. Cohen, Pearson correlation coefficient, in: Noise reduction in speech processing, Springer, 2009, pp. 1–4. [26] N. Eustathopoulos, M.G. Nicholas, B. Drevet, Wettability at high temperatures Pergamon Materials Series, Pergamon, Amsterdam, Lausanne, New York, Oxford, Shannon, Singapore, Tokyo, vol. 3, 1999. [27] R.A. McCauley, Corrosion of ceramics, Marcel Dekker, Inc., New York, 1995.
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Erosion and corrosion of alumina refractory by ingot casting steels Jens Fruhstorfer a,∗ , Leandro Schöttler b , Steffen Dudczig a , Gert Schmidt a , Patrick Gehre a , Christos G. Aneziris a a b
Institute of Ceramic, Glass and Construction Materials, TU Bergakademie Freiberg, Agricolastraße 17, 09596 Freiberg, Germany Deutsche Edelstahlwerke GmbH, Obere Kaiserstraße, 57078 Siegen, Germany
a r t i c l e
i n f o
Article history: Received 18 August 2015 Received in revised form 12 November 2015 Accepted 22 November 2015 Available online xxx Keywords: Contact angle Manganese Aluminum Silicon Oxygen
a b s t r a c t This study investigated the corrosion of alumina-based refractories by alloyed steels. The alumina crucibles were physically characterized. Four steels with varying casting temperature and alloy contents were melted under argon atmosphere and the contact angles measured. Post-mortem, X-ray diffraction, optical and scanning electron microscopy and energy-dispersive X-ray spectrometry were used to analyze the erosion and corrosion. The refractory properties did not affect erosion and corrosion. No infiltration occurred. The contents of oxygen and deoxidizing steel alloys influenced the corrosion and erosion of the matrix. Firstly, manganese reacted to manganese aluminate. Silicon followed—forming manganese alumosilicates. For highest casting temperatures (1580 ◦ C) also mullite formed. At the steel–refractoryinterface aluminum displaced the silicates, enhancing corrosion. Due to mismatching thermal expansions, spalling occurred for the steel 17CrNiMo76 respectively 1.6587 in which mullite formed. In the coarse fraction only chromium doped alumina formed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction A new generation of secondary steel metallurgy developed [1] as well as the demand for improved steel qualities is continuing. The quality requirements are met by improving the steel cleanliness by reducing non-metallic inclusions. Non-metallic inclusions are a significant problem because they act as stress raisers in the cast steels and, thus, limit the attainable mechanical properties [2,3]. During the commonly applied continuous casting of steel severe bending stresses occur at high temperatures. Steel grades which are sensitive to such stresses as well as steels for special applications are ingot cast discontinuously. These steels may be alloyed with carbon, chromium or manganese [3–7]. Ingot casting is usually performed in a bottom teeming process by pouring the steel through a refractory pipe system into the ingot mould [3,5]. According to Zhang and Thomas [2,3] non-metallic inclusions in steel are classified into indigenous and exogenous inclusions. Indigenous inclusions are deoxidation products or inclusions which precipitate during cooling. Deoxidation products can be e.g. Al2 O3 in aluminum killed steels, SiO2 in silicon killed steels or MnO in steels containing manganese [8]. Exogenous inclusions are
∗ Corresponding author. E-mail address: [email protected] (J. Fruhstorfer).
formed by chemical and mechanical interaction of the steel with its surroundings. They have, generally, the most harmful effects on the mechanical properties due to their larger defect size and surface-near location. Sources of exogenous inclusions are mainly reoxidation, slag entrainment, lining erosion and chemical reactions. Regarding the boundary interaction of an ingot cast steel with its refractory pipe system, the most serious inclusions arise from lining erosion and chemical reactions. Refractory contents with a low oxidation potential are reduced by steel elements with a high oxidation potential as e.g. Al can reduce SiO2 leading to Si and Al2 O3 [8]. Therefore, refractories containing SiO2 provide a high oxidation potential. Besides the reoxidation, compound formation is a source of additional non-metallic inclusions. In Al-killed steels containing Mn, manganese aluminates [9] and in Si-killed steels containing Mn, manganese silicates can be forming. In the complexer MnO–SiO2 –Al2 O3 system, manganese alumosilicates may be forming [2,3,10]. Also Cr-alloyed steels are reported to be highly corrosive [5–7]. Chromium oxide is forming solid solutions with alumina [11]. Furthermore, it is able to react to chromite [12,13] or alumino-chromite spinels [13]. Despite the manifold possible reactions, it was reported that high purity alumina, mullite or zirconia parts with a low amount of free silica showed a good applicability and resistance to Mnand Al-rich steels. Furthermore, they have an excellent hot erosion resistance [2,3,5].
http://dx.doi.org/10.1016/j.jeurceramsoc.2015.11.038 0955-2219/© 2015 Elsevier Ltd. All rights reserved.
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2 Table 1 Steel compositions and casting temperatures. Steel grade
1.3520 1.3816 1.4542 1.6587
Casting temp.
1522 ◦ C 1460 ◦ C 1535 ◦ C 1580 ◦ C
Table 2 Large crucible firing curve.
Chemical composition in wt% C
Si
Mn
Cr
Al
0.93 0.07 0.03 0.16
0.60 0.35 0.40 0.21
1.10 19.05 0.70 0.63
1.50 18.17 15.30 1.52
0.00 0.01 0.01 0.03
However, the corrosion and wear experiments were conducted with steels varying in their composition. Although some reactions of alloys with contents of the refractory were determined, the interaction of the alloys regarding their combined effect on the corrosion is rather unclear. Consequently, no highly corrosive steel for corrosion tests for comparing different refractory materials was determined. The purpose of this study, therefore, is to investigate the corrosion and erosion of cast alumina crucibles by four ingot casting steels of varying composition and casting temperature. A highly corrosive steel for further and future corrosion tests shall be determined. In this study, the corrosion of the alumina crucibles will be regarded after use as an indirect potential of producing inclusions and to evaluate the corrosion mechanisms [14]. Erosion and corrosion will be analyzed microscopically, chemically and mineralogically. 2. Experimental Corrosion tests with four steels varying in composition and cast alumina crucibles were carried out in an inductively heated metal casting simulator. The tested ingot casting steels were a C- and Si-rich (1.3520 resp. 100CrMn6), a Mn- and Cr-rich (1.3816 resp. X10MnCrN1818), a Crrich (1.4542 resp. X5CrNiCuNb174) steel and an Al-rich steel with a high casting temperature (1.6587 resp. 17CrNiMo76). Their chemical compositions as measured by X-ray fluorescence spectroscopy (ARL 9800, Thermo Scientific, U.S.A.) and combustion analysis (CS 244, LECO Corporation, U.S.A.) for the determination of the carbon content as well as their casting temperature are presented in Table 1. The raw materials for the refractory crucibles were alumina fractions up to 3 mm (Tabular Alumina T60/64), reactive alumina (CL370) and a re-hydratable alumina binder (Alphabond 300) obtained from Almatis GmbH, Germany. Additionally, a defoamer (Contraspum K 1012, Zschimmer & Schwarz GmbH & Co. KG, Germany), a dispersant (Castament FS 60, BASF Construction Polymers GmbH, Germany) and water were used, described in more detail by Fruhstorfer and Aneziris [15]. The composition of the castable for manufacturing large crucibles with a height of ≈25 cm and a diamter of ≈27 cm was based on a modified Andreasen model as introduced by Fruhstorfer and Aneziris [15]. In the modified model, the particle size distribution was described by a minimum and a maximum distribution modulus. In this study a maximum distribution modulus of 0.8 similar to the previous study [15] was used. In difference, a minimum modulus of 0.27 and a water content of 5.5 wt% were applied because preliminary experiments showed that with minimum moduli ≤0.27, maximum ones ≥0.8 and slightly increased water contents the flowability increased by simultaneously further decreasing the pore sizes arising during shaping. The binder content was adjusted to 3 wt%. Furthermore, 0.1 wt% of the defoamer and 0.57 wt% of the dispersant related to the solids content were added. The castables of about 21 kg mass were prepared with a compulsory mixer (R05T, Eirich, Germany). The batches were dry mixed
ϑstart in ◦ C
ϑend in ◦ C
Time in h
ϑroom 180 600 800 1600 1600 600
180 600 800 1600 1600 600 ϑroom
3 42 4 7 3 13 Free cooling
for 1 min, then the water was added and it was wet mixed for 5 min. Right after finishing mixing, the crucible moulds were filled bottom up. The moulds had a silicone covered core and were lubricated with a Teflon oil. The open top of the mould was covered with a wet tissue after complete filling. After setting for 1 d, the crucible was demoulded. The large crucibles were dried carefully due to the strong dehydration of the re-hydratable alumina binder and fired as presented in Table 2. After firing the filling volume of the crucibles was about 7.5 L. The mineralogical composition was determined by Xray diffraction (XRD) (PW1820, Philips, Netherlands) with Cu-K-␣ radiation. The corrosion experiments were conducted in an inductively heated melting unit (150 kW inductor power) of a metal casting simulator (Systec GmbH, Germany) introduced by Dudczig et al. [16,17] and Aneziris et al. [18]. The crucibles were filled with 20–25 kg steel. To reduce the oxygen content it was twice evacuated to 2.5 mbar. After the first evacuation it was flushed with argon 4.6 gas (99.996 vol% argon) to a pressure of 450 mbar and after the second evacuation to ambient pressure. Then it was inductively heated by increasing the power by 10 %/h related to the 150 kW. This resulted in an interpolated heating rate of ≈250 K/h. The steels were heated to their casting temperature (Table 1) and hold at this temperature for 1 h because in steel ingot casting the contact time before solidification is about 15–45 min in line with Zhang and Thomas [3]. By applying an hour contact time in an induction heating process the corrosion effect was maximized. Every 10 min, the temperature and oxygen content were measured with a pO2 /T-sensor-system. Cooling was carefully performed under argon atmosphere by reducing the power by 5 %/h to reduce the risk of cracking. The benefit from the large scale experiments was the resembling of the application and the obtainment of enough corroded material for detailled analysis. The crucibles were cut and analyzed after the corrosion test. The spots where mainly parallel steel flow occurred were analyzed because in the refractory pipe system also mainly parallel flow occurs. The thickness of the eroded layers was analyzed by digital optical microscopy (DOM) (VHX Digital Microscope 2000, Keyence Microscope Europe, Belgium) and the corroded layer visually or by microscopy. In the same spots on the outer area of the crucible as well as on spots of the upper crucible part, porosity and pore sizes below 150 m were measured as described in a previous study [15] with the water immersion method and mercury porosimetry, respectively. The mercury porosimetry was analyzed by the Washburn method applying an assumed contact angle of 140◦ and a mercury surface tension of 0.48 N/m. From the reacted areas, also samples for scanning electron microscopy (SEM) (ESEM XL30FEG, FEI company, Netherlands) equipped with an energy-dispersive X-ray spectrometer (EDX) (EDAX, Ametek GmbH, Germany) were prepared and analyzed. The SEM signal of back-scattered electrons was used to visualize phase differences due to their differing Z-contrasts [19,20]. To determine the mineralogy, XRD was performed. Therefore, the visually corroded sections respectively layers of the crucible were separated from each other by crushing them into pieces and classifying them according to
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3
Table 6 Pearson correlation coefficients of the mean oxygen content.
Total porosity
Open porosity
Bottom
Top
Bottom
Top
1 (1.3520) 2 (1.3816) 3 (1.4542) 4 (1.6587)
19.9 20.7 20.9 19.9
19.4 20.5 20.7 19.6
16.1 17.1 17.5 16.2
15.7 16.9 17.5 15.8
Average Total
20.4
20.1
16.7
16.5 16.6 ± 0.8
20.2 ± 0.6
Table 4 p-Values of the ANOVA for the total and open porosity (significant effects on a level ≤0.05 marked bold). Factor
Crucible number Location
p-Value of the Total porosity
Open porosity
0.0015 0.0142
0.0029 0.0955
Significant Bonferroni-corrected correlationon a level ≤ 0.01. on a level ≤ 0.05. on a level ≤ 0.1.
their apparent color. After classification, the pieces were milled to particle sizes below 63 m for the XRD measurement. Furthermore, the contact angles between cut slides of sintered parts and the different steels were measured by the sessile-drop method at the casting temperatures. A hot stage microscope as described by Aneziris and Hampel [21] was used. The casting temperature was hold for 1 h and it was continuously flushed with argon gas.
on a level > 0.1.
practical considerations in line with Schafföner and Aneziris [24] where comparable results were obtained for pressure slip castables. The pore size describing parameters were d50 = (2071 ± 111) nm and d80 = (5338 ± 1675) nm in the pore size range <150 m. The di means that i % of the pores were smaller than this size. The samples were taken from the bottom location. In dependence on the crucible resp. the porosities, no significant differences of the pore sizes were detected. In comparison to the data of Fruhstorfer and Aneziris [15] (d80 = 3691 nm) the d80 pore size tended to increase slightly. It was found in the previous study that for denser castables the pore sizes shifted towards smaller values, increasing their di in the pore size range <150 m. Consequently, the overall pore sizes were further shifted towards smaller values by using slightly more fines and 0.5 wt% more water.
3. Results and discussion Firstly, the refractories and, afterwards, the melting tests were investigated. Finally, the corrosion results were correlated with the refractory and steel properties. 3.1. Alumina refractory characterization
3.2. Melting tests
The qualitative XRD analysis of the crucibles detected only corundum and traces of -alumina after firing. The total and open porosities of the crucibles are listed in Table 3. The total porosity was (20.20 ± 0.59) vol%. In a previous study [15], 5 wt% water were added to a self-flowing castable resulting in an average total porosity of 19.4 vol% and 15.7 vol% open porosity. By using 5.5 wt% water in the current investigation, each porosity increased for about 1 vol%. To investigate whether the hydrostatic pressure causes a density gradient, an analysis of variances (ANOVA) was used to determine the main effects of the factors ‘Crucible number’ and ‘Location’ (bottom or top) on the total and open porosity. An ANOVA tests means on their statistical difference [22]. The tests were performed with the statistics software R [23]. An effect was regarded significant when its p-value was ≤0.05. Table 4 presents the resulting p-values. As the factor ‘Location’ had a significant influence on the total porosity, in the cast material a density gradient arose. However, the influence of the ‘Crucible number’ was more significant. Thus, the statistical differences of the factor ‘Location’ were negligible for
During the holding time, every 10 min the temperature and oxygen content were logged. The averages and standard deviations are presented in Table 5. The specified casting temperatures were met satisfactorily. The highest inaccuracy showed the melting test of the Mn- and Cr-rich steel 1.3816—the average temperature was about 10 K below the required one. The arising oxygen content was correlated with the average melting temperature as well as with the steel and crucible properties. Table 6 presents the Pearson product–moment correlation coefficients [25] and the p-values of the t-tests on difference from zero [22,23]. In the present study, coefficients below 0.5 are regarded to indicate a poor linearity, coefficients between 0.5 and 0.8 a medium one and larger coefficients a strong linear correlation. It can be seen in Table 6 that the crucible properties (porosities and pore sizes) showed no significant correlations with the oxygen content. The determined significant differences between the crucibles (see Table 4 in the prior subsection) had no practical influence
Table 5 Logged data during steel melting in the large crucibles (* specified by steel manufacturer). Steel grade
Casting temp.*
Logged temperature in ◦ C
1.3520 1.3816 1.4542 1.6587
1522 ◦ C 1460 ◦ C 1535 ◦ C 1580 ◦ C
1519 1451 1536 1580
± ± ± ±
9 26 6 10
Logged oxygen content in ppm 10 1 20 36
± ± ± ±
1 2 1 2
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Fig. 1. Dependency of the oxygen content in the melt on the highly significant correlation properties.
on the oxygen content. But the temperature showed a significant correlation with a strong correlation coefficient indicating a high linearity of the relation (see Fig. 1(a)) similar to the observations of Kim and McLean [9]. Manganese, silicon and chromium are possible deoxidants [8] and, therefore, their correlation coefficients were negative and significant. The correlations were medium because the relations were supposedly not linear due to an exponential nature or occurring interactions—in the present study the relation on the manganese content was exponential as can be seen in Fig. 1(b) similar to the results of Kim and McLean [9]. The dependency of the oxygen content in the melt on the aluminum content was positive and significant although Al is a known strong deoxidant. Therefore, the observed relation is presumably indirectly determined. In Fig. 1(c) can be seen that at a level of 0.01 wt% Al two levels of resulting oxygen contents (≈3 ppm and ≈20 ppm) arose. The lower oxygen content resulted when the Mnand Cr-rich steel 1.3816 was melted at ≈1451 ◦ C—Al acted deoxidizing. The higher oxygen content arose during melting of the Cr-rich steel 1.4542 at ≈1536 ◦ C. The aluminum amount interacted with the temperature. For higher temperatures Al presumably affected the corrosion significantly and, therefore, did not show a deoxidizing behavior. 3.3. Corrosion analysis Wetting, erosion and corrosion were analyzed. Firstly, the wetting was analyzed by regarding the contact angles. An ANOVA showed statistically that the contact angle did
not change during 1 h holding time. Therefore, the angles at 0 and 60 min were used for the analysis. The values are presented in Table 7. Contact angles above 90◦ characterize a non-wetting behavior and contact angles above 110◦ a non-reactive one according to Eustathopoulos et al. [26]. It was, therefore, tested with one sample Bonferroni-corrected t-tests if the measured contact angles were greater than 90◦ and lower than 110◦ . Only the Mn- and Cr-rich steel 1.3816 wetted the alumina ( < 90◦ ) and only the C- and Si-rich steel 1.3520 showed a non-reactive behavior ( ≥ 110◦ ). The correlations of the contact angle with the steel, crucible and melting properties are presented in Table 8. On a significance level of 0.05 after Bonferroni correction, only the manganese content and the d80 pore size correlated significantly negative—leading to decreased contact angles. In line, an increasing manganese content of iron–manganese alloys on solid alumina under argon atmosphere at 1560 ◦ C led to decreasing contact angles [10]. An increasing d80 means that the sample contained less small pores which contribute to the surface area. And a decreasing surface area commonly contributes to an increasing advancing contact angle [26]. The corrosion was firstly analyzed macroscopic, regarding the depth of the eroded and corroded section including the eroded part. Tables 7 and 8 present the values and the correlation coefficients. The eroded section correlated on an adjusted significance level of 0.05 with no measured property. Highly probable interactions between the properties on the erosion could cause this non-significance. Therefore, also the coefficients significant on a
Table 7 Large crucible corrosion test results.
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Table 8 Pearson correlation coefficients.
Fig. 2. SEM image of alumina corroded by C- and Si-rich steel 1.3520.
Significant Bonferroni-corrected correlation. on a level ≤ 0.01. on a level ≤ 0.05. on a level ≤ 0.1. on a level > 0.1.
level of 0.1 were marked in Table 8. The erosion, thus, correlated with the oxgen and aluminum content. Both correlations showed a medium linearity affirming an interactive behavior. An increasing oxygen content of the melt led to increased erosion in line to the reviewed results by Zhang and Thomas [3] where the fatigue lifetime of bearing steels decreased exponentially with increasing oxygen contents. For every corrosive attack of liquid steel, oxygen was a reactant—independent of the other reactants. Furthermore, increasing amounts of Al thickened the eroded layer. Supposedly, after the first reactions between the crucibles and the other elements, the Al reacted with the first corrosion product—decomposing it or forming a more complex product. Both iterated corrosion mechanisms would explain the non-deoxidizing behavior observed in the prior subsection. The complete corrosion correlated with the silicon content negatively on a significance level of 0.1 in addition to the oxygen and aluminum content. As silicon acted deoxidizing (see prior subsection), it reduced the oxygen amounts and, consequently, the corrosion. But the other deoxidants like Mn did not show significant influences, although their oxidation potential is higher than the one of silicon [8]. Possibly, the other deoxidants reacted with the refractory in elementary or oxidic form. Oxidized silicon has also a high potential to form crystalline or amorphous silicate structures with manganese oxide or alumina. Anyway, such subsequent reactions would not necessarily increase the extent of the complete corrosion. Regarding the aluminum effect, the presumption from the prior subsection that it strongly affects the corrosion was confirmed. The XRD results corresponded to the colors of the layers. In the crucible with no visible layer, corroded by the C- and Si-rich steel 1.3520 only alumina was detected by XRD. The non-reactive behavior confirmed the wetting result. The corroded layer of the Mn- and Cr-rich steel 1.3816 had a tinted yellow color, corresponding to manganese aluminate. Additionally, some amorphous parts existed. The amorphous parts were qualitatively evaluated by regarding the peak heights of the chromium doped alumina in combination with the following presented EDX measurements. In case of the Cr-rich steel 1.4542, the corroded layer was parted into an inner (near steel) red–violet layer of about 1 mm thickness and
a subsequent tinted yellow layer. The red–violet layer contained mainly chromium doped alumina and traces of compounds of the system Mn–Si–Cr–O. It presented also the largest amorphous parts. Of the tinted yellow layer not enough material for a separate XRD analysis was obtained. Anyway, as for the complete corroded section also manganese aluminate traces were found which were not detected in the red–violet layer, the tinted yellow layer contained manganese aluminate. The eroded respectively partly spalled layer of the Al-rich steel 1.6587 with the highest casting temperature was tinted red colored, corresponding to chromium doped alumina and few manganese aluminate with some amorphous parts. The subsequent corroded layer was again tinted yellow due to manganese aluminate but had in contrast to the Mn- and Cr-rich steel 1.3816 very low amorphous parts. The XRD study of the individual layers returned also that mullite formed in the tinted yellow section of the alumina corroded by 1.6587. Between the mullite containing tinted yellow layer and the prior tinted red one containing chromium doped alumina, a thermal expansion mismatch existed. This mismatch led to spalling respectively delamination and explained the strongly increased erosion (see Table 7). This is in line with a study reviewed by McCauley [27] of a SiC–Al2 O3 composite, where upon oxidation at 1530 ◦ C a zone containing mullite and amorphous alumo-silicates formed. Delamination was reported as a possible consequence. To determine the corrosion reactions in more detail, SEM and EDX measurements were analyzed. The results of the EDX measurements are presented in Table 9. The SEM image in Fig. 2 presents a section of the alumina crucible corroded by the C- and Si-rich steel 1.3520. Although by XRD only alumina was detected, in the image additional phases were observed. The bright areas contained mainly iron oxide, which was found to a depth of about 40m. The gray areas corresponded to manganese aluminate found to a depth of about 30m. Down to a depth of about about 90m rarely iron aluminates were found. The porosity was not filled with steel—infiltration did not occur. The SEM image in Fig. 3(a) presents a section of the alumina refractory corroded by the Mn- and Cr-rich steel 1.3816. The XRD results showed, that only manganese aluminate was detected. The EDX measurements presented in Table 9 were in line and visualized for the matrix (finest particles) in Fig. 3(b) dependent on the depth. A measurement on a coarse grain showed that it reacted only to a small extent. Only to a depth of about 300m, additionally, chromium resp. chromium oxide doping was detected in amounts below 0.5 mol%. With increasing depth, the content of manganese oxide of the manganese aluminate compositions decreased—the non-stoichiometry explains the increased amorphous parts observed by XRD. Additionally, it can be said that the corrosion rate was controlled by diffusion. The aluminate compositions were detected to a depth of 1.4 mm. Silicon contents in the aluminate were detected to a depth of 1 mm. Supposedly, it
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followed after the manganese reacted with the refractory. It appeared in the non-stoichiometric and presumably amorphous aluminate structures and hence was not detected by XRD. This affirmed the presumption of the iterated corrosion. The refractory was not infiltrated by the steel. In Fig. 4 the SEM image of an upper section of an alumina sample corroded by the Cr-rich steel 1.4542 is presented as well as EDX measurements dependent on the depth. At the interface to the steel, only chromium doped alumina was found to a depth of about 100 m. Subsequently, a layer of manganese alumo-silicates followed to a depth of 200m while larger particles corresponded to chromium doped alumina. In the following section, chromium doped manganese aluminates were found besides chromium doped manganese alumosilicates. The chromium contents decreased with increasing depth till 800m. For depths below ≈2.1 mm only manganese aluminates were found—from there also the tinted yellow layer started. The red–violet layer formed due to the corrosion. Although, many different compositions were detected by XRD, from the EDX images is clear that also many amorphous non-stoichiometric structures formed. No infiltration was observed. Firstly, it seems, the manganese reacted with the alumina refractory. The increased casting temperature led to an increased crystallization of the compound compared to the manganese aluminate formed during corrosion by the Mn- and Cr-rich steel 1.3816. Subsequently, the silicon reacted with the refractory and the manganese aluminate, forming silicates besides the manganese aluminate. Towards the interface to the steel, the aluminate was gradually displaced by silicon containing reaction products. Anyway, at the interface to the steel only chromium doped alumina was detected. Consequently, the silicates were displaced at the interface by the aluminum at higher temperatures. Thus, the aluminum decomposed the previously formed corrosion products. The freed educts like manganese oxide and silica were hence available for new corrosion reactions. The aluminum enhanced the corrosion
Fig. 3. Alumina corroded by Mn- and Cr-rich steel 1.3816.
Table 9 Average EDX results for different depths and probable corresponding compounds Steel grade
1.3520
1.3816
1.4542
1.6587
Affected particles
Depth in m
Finest Fine Finest Finest Coarse Finest Finest Finest Finest Fine Finest Finest Finest Fine Finest Finest Finest Finest Finest Finest Finest Coarse Fine Finest Finest Finest Finest Finest Finest
20 20 80 300 300 700 1000 1400 70 150 160 310 310 550 560 800 1300 2050 2050 3200 50 50 350 525 700 1400 2500 5000 6300
Chemical composition in mol% Al2 O3
MnO
MgO
SiO2
Fe2 O3
Na2 O
Cr2 O3
Mo2 O3
CaO
52.2 2.1 61.5 49.9 96.6 71.1 74.2 95.7 93.5 94.8 43.0 33.2 54.0 94.9 46.6 42.3 38.8 56.9 40.2 59.2 34.7 98.7 48.0 17.5 17.9 53.3 52.3 60.0 52.1
41.0 0 0 48.6 1.3 17.3 10.1 1.5 0 0 28.1 44.7 7.9 0 42.7 10.0 10.0 36.2 15.3 34.7 7.7 0 36.8 16.9 17.7 41.7 43.8 34.5 44.5
2.0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 3.9 0 3.1 0 0 10.3 0.9 0.9 1.8 1.4 0.4 0.8
0.9 0 0 1.1 2.1 5.8 8.3 1.6 0 0 14.1 0.5 33.1 0 0.9 40.6 43.4 0 37.1 0 46.6 0 0 51.2 50.9 0 0 2.1 0
2.9 91.8 24.2 0 0 0 0 0 0 0 0 0 0 0 1.5 0.3 0.3 3.1 0.6 3.0 1.0 0 0.9 0 0 2.3 2.3 1.7 2.6
0.5 3.2 7.8 0 0 5.7 6.8 1.2 0 0 0.8 0.7 3.3 0 0 6.4 7.6 0 4.4 0 4.9 0 0 7.4 7.7 0.5 0 1.4 0
0.3 0.3 0 0.4 0 0 0 0 6.5 5.2 14.0 20.9 1.7 5.1 8.4 0.5 0 0 0 0 0.2 1.3 3.7 0 0 0.6 0.2 0 0
0 2.5 1.2 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0
0 0 0 0 0 0.2 0.6 0 0 0 0 0 0 0 0 0 0 0 2.3 0 4.0 0 0 5.0 4.2 0 0 0 0
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Fig. 4. Alumina corroded by Cr-rich steel 1.4542.
therefore indirectly, which explained finally the correlations and confirmed the theory of the iterated corrosion. Fig. 5 shows the upper section of the alumina corroded by the Al-rich steel 1.6587 with the highest casting temperature and EDX measurements. The white ball-shaped structures were mainly iron (e.g. 99.1 at% Fe, 0.9 at% Ni). As no infiltration was observed, the iron balls were formed most likely by a reduction of iron compounds. In the SEM-Fig. 5(a) can be seen that the surface was eroded strongly for at least 200 m besides the observed spalling reported in Table 7. This is in the size range of the Cr-doped alumina layer
Fig. 5. Alumina corroded by Al-rich steel 1.6587 with the highest casting temperature.
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observed for the corrosion by the Cr-rich steel 1.4542. Presumably due to the erosion, the alumina layer was not found by EDX in Fig. 5(c). That it existed also here was supported by the increasing alumina contents for depths more near to the surface. The increased hot surface erosion might be caused by the high treatment temperature of 1580 ◦ C. In the coarse fraction at the surface-near locations Cr-doped alumina and in the fine but not finest fractions, manganese aluminate was found. While the chromium was solid soluted in alumina, this confirms the hypothesis of the iterated corrosion, starting with the manganese and following silicon because in the finest fraction to a depth of about 1 mm manganese alumino-silicate compositions were found. As they were not detected by XRD they were presumably amorphous. Only to a depth of 50m the silicates also contained chromium. The following layer was the tinted yellow one corresponding to mullite and manganese aluminate by XRD and chromium doped alumina in the coarse fraction. The 2:1-mullite layer formed in this sample between 1 and 1.5 mm depth because there the necessary silicon amounts were reached as visualized in Fig. 5(c). It was, therefore, situated behind the spalling region, confirming that in interaction with the temperature the mullite formation and the thermal expansion mismatch led to the spalling. Below a depth of 1.5 mm, mainly manganese aluminate spinels containing chromium were detected. Chromium was detected to a depth of 2.5 mm. The subsequent manganese aluminate spinel section reached to a depth of 6.3 mm. In contrast to the EDX results of the Mn- and Cr-rich steel 1.3816, no gradually decreasing manganese content was observed, explaining the higher crystallinity. Furthermore, it can be said that the corrosion rate by the steels 1.6587 and 1.4542 with higher casting temperatures was chemical reaction controlled. The chromium doping interacted with the temperature. For larger amounts and higher temperatures, the depth the chromium reached increased. Anyway, it seemed not to cause harmful damage to the refractory. The result that the Al-rich steel 1.6587 with the highest casting temperature led to the most harmful erosion and corrosion was conform to practical observations in Deutsche Edelstahlwerke GmbH steelworks. 4. Conclusion The erosion and corrosion of alumina crucibles by ingot casting steels of varying composition was investigated. • The contents of oxygen and the deoxidizing steel alloys significantly influenced erosion and corrosion. • An iterated corrosion occurred dependent on the temperature. The corrosion started with manganese reacting to manganese aluminate with the alumina refractory. Subsequently, a transition to manganese alumo-silicates occurred with silicon, firstly occurring besides the aluminates. At the interface, aluminum displaced these silicates in a sharp transition. Thus, the silica was again available for further corrosive reactions. Consequently, the aluminum content enhanced the corrosion strongly. • The highly corrosive steel 1.6587 (17CrNiMo76) led to the largest erosion and corrosion of the alumina-based refractories. Thus, it is recommended for further corrosion tests for comparing different refractory materials. Only during this corrosion test a mullite containing zone with a different thermal expansion formed due to the high casting temperature of 1580 ◦ C, leading to spalling. The high casting temperature and its interactions with the steel alloys, consequently, made this steel the most corrosive.
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Please cite this article in press as: J. Fruhstorfer, et al., Erosion and corrosion of alumina refractory by ingot casting steels, J Eur Ceram Soc (2015), http://dx.doi.org/10.1016/j.jeurceramsoc.2015.11.038