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Effect of titanium on the as-cast microstructure of a 16%chromium white iron
Materials Science and Engineering A 398 (2005) 297–308
Effect of titanium on the as-cast microstructure of a 16%chromium white iron A. Bedolla-Jacuinde ∗ , R. Correa, J.G. Quezada, C. Maldonado Instituto de Investigaciones Metal´urgicas, Edificio “U” Ciudad Universitaria, Universidad Michoacana de San Nicol´as de Hidalgo, Morelia Michoac´an, C.P. 58000, Mexico Received 17 August 2004; received in revised form 15 March 2005; accepted 29 March 2005
Abstract This research work studies the effect of systematic titanium additions (up to 2 wt.%) to a 16%Cr, 2.5%C white cast iron. The study was undertaken in six laboratory made alloys with different titanium amounts. Alloys were melted in an open induction furnace by using high purity raw materials. Such additions caused small hard titanium carbide particles to precipitate within the proeutectic austenite therefore promoting a strengthening of matrix; such particles also contributed to increase bulk hardness of the overall alloy. A structure refinement, as measured by the secondary dendrite arm spacing, was also observed as the titanium amount was increased. Titanium carbide precipitation caused a small decrease in the eutectic carbide volume fraction. The fracture toughness remained constant since the strengthening of matrix was compensated with a decrease in the volume of eutectic carbides. According to this study, titanium can be used as an alloying element to increase the hardness and perhaps wear resistance without affecting fracture toughness in high-chromium cast alloys. The results are discussed in terms of the precipitation nature of such small hard titanium particles. © 2005 Elsevier B.V. All rights reserved. Keywords: Cast iron; Titanium; Microstructure; High-chromium
1. Introduction High-chromium white irons can be described as in situ composites with large, hard eutectic and/or proeutectic M7 C3 carbides in a softer iron matrix (commonly austenitic in the as-cast conditions and mostly martensitic after a destabilization heat treatment). These microstructure combinations make high-chromium white irons very suitable for applications where abrasion resistance is the main requirement. Their exceptional abrasive and erosive wear resistance results primarily from their high volume fraction of hard carbides, although the toughness of the matrix also contributes to the wear resistance [1]. Attempts to improve the wear resistance in highchromium irons have led researchers to try different alloying elements such as vanadium, tungsten, titanium and niobium [2–14]. Even the addition of carbides such as TiC [15] has ∗
Corresponding author. E-mail address: [email protected] (A. Bedolla-Jacuinde).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.072
been undertaken to obtain composites in these irons. The aim of these additions is usually to achieve some modification of the eutectic carbide structure by obtaining harder carbides, though they may improve the hardenability of the matrix since these carbide forming elements may partition also to the matrix. Laird [16] and Ono et al. [17] have widely studied the partition coefficient or segregation ratio for several elements in high-chromium irons. It has been recognised that a possible strategy for improving the toughness of white iron alloys as well as the wear resistance under sliding conditions involves the refinement of the eutectic carbide structure by producing finer, more globular carbides [18–20]. Rapid cooling [21,22] or lower superheat [16] have been used as a means of increasing nucleation and hindering carbide growth, producing finer carbides. Conversely, slow cooling of castings results in larger dendrite arm spacing, reducing the number of sites for nucleation of the eutectic and hence, larger, coarser, eutectic forms [23,24]. Alloying additions of some elements that segregate to the matrix/carbide interphase, such as rare earths, silicon,
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boron, etc., have also been used to modify the eutectic carbide structure [25–34]. Generally, these attempts to modify the eutectic carbide structure through alloying elements have had limited success. Wear resistance may be increased by increasing the eutectic carbide volume fraction in the alloy but at the cost of further decreasing fracture toughness. Researchers, therefore, have the challenge of improving wear resistance without reducing fracture toughness. The high amount of eutectic carbides play a critical role in the fracture behavior of white iron alloys, and the toughness has been seen to deteriorate as the proportion of carbides in the structure increases [35]. Since a high proportion of carbides in the structure is required to maximise the wear resistance, many researchers have sought to improve the toughness through modifications to the carbide structure [26,32]. The carbide attributes which have been manipulated in an attempt to improve toughness include volume fraction [35], carbide size [36,37], and inter-carbide spacing, and carbide morphology and shape [36]. These studies have found, with varying degrees of success, that the toughness of high-chromium cast irons may be improved by decreasing the proportion of carbides in the structure, decreasing the carbide size, increasing the inter-carbide spacing, and globulizating the carbide shape. From this study it is expected to improve wear resistance in a 16%Cr white cast iron by adding titanium as an alloying element. Several carbide forming elements have been used for this purpose, as mentioned above, but the use of titanium has been very limited due to the difficulties found during ironmaking. Titanium is a very oxidation prone element and special conditions for alloying are required when using an open induction furnace (not controlled atmosphere). The hypotheses for using titanium to improve wear resistance without decreasing fracture toughness are: • Titanium is a strong carbide forming element; since TiC has a high formation temperature, it will be the first phase precipitating during solidification. Therefore, a proeutectic matrix reinforced with hard small TiC particles is expected at the end of the solidification process. The reinforced matrix, in turn, is expected to improve wear performance. • It is also expected that TiC formation refines the structure since such particles may act as nuclei for the austenite phase. • As a result of this, a fine microstructure with isolated hard small TiC within the matrix would be expected which improves wear performance.
2. Experimental procedure The alloys used in the present study were made in a laboratory induction furnace by using high purity raw materials. Thirty kilograms of the alloy were melted and poured at 1450 ◦ C into 5 kg capacity SiC crucibles where the titanium
additions were undertaken. Therefore, six crucibles with different additions of titanium were obtained which were then poured into sand moulds to obtain square 2.5 cm × 2.5 cm cross section cast bars. The amount of titanium was zero for the first ingot, 0.1% for the second, 0.7% for the third, 1.5% for the fourth, 1.7% for the fifth, and 2% for the sixth. Chemical analysis was undertaken by spectrometry from chill samples obtained during casting each ingot, and the results are shown in the next section. Samples for metallography, hardness, wear, and fracture toughness tests were cut from the cast bars. Sample cutting was carried out in a Discotom by using an abrasive wheel of alumina; cutting was done as slow as possible in order to avoid excessive overheating that may cause cracking in the samples. Additionally, copious amounts of water as coolant were used. Material’s characterization was undertaken by light microscopy, scanning electron microscopy (SEM), transmission electron microscopy (TEM), microanalysis by energy-dispersive spectrometry (EDS), X-ray diffraction (XRD) and PC based image analysis. Samples for metallography were prepared in the traditional way; roughed by using abrasive paper and then polished on nylon cloths by diamond paste. Once polished, the specimens were etched with one of the two different reagents: Villela’s reagent (5 ml HCl, 1 g picric acid in 100 ml ethanol) for 30 s to reveal the microstructure; or in a solution of 50 ml FeCl3 plus 20 ml HCl in 930 ml ethanol for about 3 h for deep etching. The later reagent removes part of the matrix without affecting the carbide phase providing a very high contrast between matrix and carbides when the structure is observed under the light microscope. Carbide volume fraction and carbide size were measured by image analysis using digital images taken on the light microscope at about 250×. For this purpose, samples were deep-etched to get high contrast. Twenty digital micrographs were processed using the software Sigmascan V5. Further characterization was undertaken by XRD in a SIEMENS 5000 diffractometer using Co K␣ radiation in a 2θ range of 30–130◦ . A JEOL 6400 SEM was also used for imaging and microanalysis. In addition, some thin foils for transmission electron microscopy were prepared from the iron with higher titanium amount. Thin sections were cut in discs of 3 mm and milled to 80 m, then dimpled to 15 m and finally ion milled to perforation. Samples were analyzed in a Philips Tecnai F20 TEM. Microhardness of the matrix in the different irons was measured on metallographic samples lightly etched with Vilella’s reagent. About 15 measurements for each sample were undertaken using a microhardness tester LECO M-40 with Vickers diamond indenter applying a load of 100 g for 15 s. Hardness, on the other hand, was measured on the Rockwell C scale for the bulk material. Between 10 and 15 measurements were undertaken on as-polished samples using a Rockwell hardness tester with 150 kg load and a diamond pyramid indenter typical for this scale.
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Table 1 Chemical composition of the high-chromium cast irons used for the analysis Element (wt.%)
Sample 1
Sample 2
Sample 3
Sample 4
Sample 5
Sample 6
C Cr Mo Ni Mn Si Ti
2.40 15.92 2.91 2.52 1.77 1.41 0.0
2.46 16.02 2.88 2.62 1.78 1.48 0. 11
2.48 16.02 2.92 2.49 1.90 1.4 0.72
2.45 15.82 2.95 2.53 1.86 1.45 1.53
2.61 16.12 2.90 2.55 1.84 1.45 1.68
2.58 15.90 2.89 2.60 1.77 1.45 2.02
Fracture toughness testing was carried out by three point bending in a Hounsfield universal testing machine. The samples whose dimensions were 2.5 cm × 2.5 cm × 15 cm, were U-notched and tested at a speed of 0.5 mm/min according to the ASTM E 399-90 standard [38] but without pre-cracking. The reported KIC values are the mean of three tests.
3. Results and discussion Table 1 shows the chemical composition for the six white irons used for the study. Note the similar amounts of carbon, chromium, nickel, molybdenum, manganese and silicon obtained for the alloys; the only varied element was titanium from 0 to 2.02%. Since the pouring temperature and cooling rate were the same for all the irons, variations in microstructure as well as in mechanical properties can be attributed to the effect of different titanium amounts. 3.1. Microstructure Fig. 1 shows an optical micrograph where the general microstructure of the iron with no titanium additions can be observed. The microstructure consists of hard eutectic carbides in a soft austenitic matrix. The solidification path fol-
lowed by this kind of alloys has been clearly summarized by Tabrett et al. [39]; being a hypoeutectic alloy, solidification starts with the formation of austenite dendrites which grow while the temperature decreases and the remaining liquid reaches the eutectic composition; then the eutectic reaction takes place and the coupled austenite/carbide eutectic is developed. Austenite remains as a metastable phase at room temperature due to the high amount of alloying elements, particularly molybdenum and nickel that increase hardenability and carbon that lowers the martensite start temperature MS at temperatures below room temperature [1]. XRD studies on this iron also show the presence of small amounts of martensite and the molybdenum rich M2 C carbide in addition to austenite and the eutectic M7 C3 carbide, Fig. 2. The presence of M2 C has been reported by several authors [1,16,17,25,40] in these alloys. According to them, most of molybdenum in the alloy partitions to the eutectic M7 C3 carbide and some dissolves in austenite (the Mo dissolved in austenite is the one that really helps to improve hardenability in the alloy), but a small part of Mo also forms a type M2 C carbide due to the presence of high amounts of carbon in the iron. Such a carbide forms as a eutectic and is the last part to solidify. It is commonly found joined together with the M7 C3 but it can be clearly distinguished from it from the lamellar shape of the former. Fig. 3 shows an SEM micrograph indicating the presence of the M2 C carbide and Fig. 4 shows detail of such a carbide in a deep-etched sample along with its composition, which clearly indicates is a molybdenum-rich carbide. The presence of martensite in these alloys has also been widely reported elsewhere [5,14,16,25]. During eutectic solidification the M7 C3 carbide which grows along the austenite, absorbs carbon from its surroundings and a narrow area at the austenite/carbide interphase becomes impoverished in carbon. The lack of carbon in these zones of austenite increases the MS temperature which allows these areas of austenite to transform to martensite during subsequent cooling down. Therefore it is common to find eutectic M7 C3 carbides surrounded by laths of martensite. Fig. 5 shows an SEM micrograph to confirm this phenomenon. 3.2. Effect of titanium on the microstructure
Fig. 1. Light micrograph showing the general microstructure of the iron without titanium additions. The structure is basically composed of eutectic M7 C3 carbides in an austenitic matrix. Vilella’s reagent for 30 s.
The microstructure of the irons with titanium additions is similar to the base iron, the only difference is the presence of titanium carbide (or carbonitride) mostly within the
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Fig. 2. XRD traces for the iron with no titanium additions showing the microstructure to consist of austenite, M7 C3 and small amounts of martensite and M2 C.
Fig. 3. SEM micrograph showing the presence of lamellar M2 C molybdenum-rich carbide.
Fig. 5. SEM micrograph of a high-chromium cast iron showing the presence of martensite laths at the carbide/matrix interface due to carbon depletion at these areas.
Fig. 4. SEM micrograph of a deep-etched sample showing detail of the M2 C carbide and the EDS analysis evidencing it is a molybdenum-rich carbide, the major element that determines the stoichiometry.
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Fig. 6. Sequence of SEM micrographs showing the increase in the amount of TiC (arrowed) within the austenitic matrix as the titanium content increased in the alloy.
proeutectic austenite matrix. Fig. 6 shows the microstructural evolution of the irons as the amount of titanium increases in the alloys. Note the increase in the number of small titanium carbides in the proeutectic austenitic matrix. XRD traces on the 2%Ti iron shows a couple of peaks indicating the presence of TiC (Fig. 7). Furthermore, chemical composition of these particles was clarified by EDS in the SEM (Fig. 8) and results confirm the particles to be TiC. In addition of titanium and carbon, some other elements such as molybdenum, chromium and iron are present in the particle, according to
the EDS. A more detailed characterization of these particles was undertaken by TEM as shown from Fig. 9. That figure shows a bright field micrograph where a hexagonal-shaped TiC particle can be observed, which was oriented to a zone axis [0 1 1] perpendicular to the plane; previous TEM calibration indicated good correlation between the measured distances from the transmitted beam to the indexed spots and the ratios obtained from databases. In addition, EDS analysis by a fine electron beam indicates only the presence of carbon and titanium in the particle. This later observation confirms
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Fig. 7. XRD traces for the iron with 2.02% titanium additions showing the microstructure to consist in austenite, M7 C3 , small amounts of martensite, M2 C and TiC.
that the presence of any other element in the EDS analysis undertaken by SEM might be due to the matrix interaction with the beam. Fig. 10 shows a bright field TEM micrograph where a TiC particle appears white and the dark austenitic ¯ as shown by the matrix is oriented with the zone axis [1 1¯ 2] diffraction pattern. In addition, EDS on the matrix indicates the presence of elements such as chromium, molybdenum, nickel, manganese, silicon, carbon and titanium. Cr, Mo, Fe, C and Mn also may be found in the carbide phase since they partition to both carbide and matrix phases [14,17,23,39,41]. From the present study, the specific temperature at which these fine TiC particles (1 m size or less) formed was not determined. However, it may be deduced that these fine particles formed at some temperature higher than that of austenite dendrite formation (when the alloy is still liquid) and that these small carbides were entrapped by the growing austenite dendrites. Since these TiC particles are the first to solidify, it is likely they serve as nuclei for proeutectic austenite. Actually, a refinement of the secondary dendrite arm spacing from 25
to 18 m was observed when titanium was increased from 0 to 2.02% (Fig. 11). Such a refinement has also been observed and reported in a previous work [14] and is attributed to the precipitation of TiC particles when the alloy is still liquid. If these particles actually function as nuclei for austenite, there must be a crystallographic relationship between some planes of TiC and austenite; however, such a relationship was not studied in the present work. Fig. 12 shows an SEM micrograph of the 2%Ti iron where a uniform particle distribution within the matrix is observed. However, such a distribution was not observed all around the sample. Some areas where agglomeration of TiC particles occurred were also observed in the samples with high Ti additions (see Fig. 13). Fig. 14 shows a mosaic of dark field micrographs of agglomerated TiC particles where austenite appears dark and TiCs are white (the image was generated by the diffracted [2 0 0] spot of TiC indicated in the diffraction pattern). Also a diffraction pattern of the area in the mosaic was undertaken, which shows spots of austenite and
Fig. 8. SEM micrograph showing two TiC particles in austenite and the EDS showing the chemical composition of the arrowed particle. 2%Ti iron sample, Vilella reagent, 45 s.
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Fig. 9. Bright field TEM micrograph showing a hexagonal-shaped TiC particle (a) with its corresponding selected area diffraction pattern (b) and chemical composition as undertaken by EDS (c). Sample of the 2%Ti iron.
rings formed for the several TiC particles in the area. Such rings are not continuous due to the large size of particle and perhaps some degree of preferred orientation. This last observation reinforces the theory that there must be some crystallographic relationship between austenite and TiC, so that austenite could precipitate on the formed TiC. However, not all the particles whose size is about 1 m were well dispersed in the matrix; Fig. 15 shows SEM micrographs where coarser (up to 4 m size) TiC appeared attached to the eutectic carbides at the M7 C3 /austenite interface. Such a phenomenon may be explained as follows: a large number of TiC particles precipitate at high temperatures and most of them serve, perhaps, as nuclei for the subsequent austenite
formation. Once primary austenite dendrites have formed, they will grow rejecting some alloying elements into the liquid and collide with some other TiC particles. Since titanium has a low partitioning coefficient to austenite [42], a titanium rich layer is expected at the periphery of the austenite arms. That increase in solute probably induces a constitutional supercooling which facilitates coarsening of particles. The carbon needed for the carbide formation is easily removed from the melt on account of its high diffusivity. Thus, whenever the free carbide particles collide with the interface, they are pushed away by the dense layer. Sometimes, these particles are caught by the faster moving interface and appear at the centre of the dendrites with smaller sizes; they may have been
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Fig. 10. Bright field TEM micrograph showing a hexagonal-shaped TiC particle in matrix (a), (b) selected area diffraction pattern of matrix showing to be austenite, and (c) the chemical composition of matrix by EDS showing the presence of dissolved elements such as Cr, Ni, Mo, C, Mn, Si and even Ti. Sample of the 2%Ti iron.
entrapped at the early stages of their coarsening or they nucleated austenite. Kesri and Durand-Charre [7] have reported similar observation for the NbC formation during solidification of Fe–Cr–C alloys. Fig. 16 shows variations in the eutectic carbide volume fraction as the amount of titanium increased in the irons. It is well known that in iron–carbon alloys, particularly in
cast irons, high solidification rates favor the formation of the metastable carbide/austenite eutectic instead of the stable austenite/graphite eutectic independently of the presence of other graphitizing or carbide forming elements. For the present study, all the alloys were solidified at the same rate and the differences observed are attributed to the effect of titanium. Such differences are due to the precipitation of TiC
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prior to the eutectic solidification. The formation of these carbides consumes part of the carbon present in the alloy so that during eutectic solidification the volume of M7 C3 is diminished. These results are in agreement with some works where titanium and/or niobium have been used as alloying elements in high-chromium irons [3,7,11,14,41]. 3.3. Effect of titanium on the mechanical properties
Fig. 11. Secondary dendrite arm spacing as a function of the titanium content in the alloy.
Fig. 12. SEM micrograph of the 2%Ti white iron showing the well distributed small TiC particles in the austenitic matrix.
Fig. 17 shows the increase of matrix hardness as the amount of titanium increases in the alloy. This logical phenomenon is achieved due to the reinforcement of austenite by a large number of hard small well dispersed TiC particles. It is also possible that titanium dissolved in matrix for the irons with higher Ti amounts (see Fig. 10), may have contributed to the matrix strengthening by solid solution; however, such a contribution would be negligible compared to the dispersion strengthening. For this case, a considerably increase in hardness from 330 to 500 Vickers100 was observed. The maximum volume of TiC was ∼3% for the iron with 2%Ti with particles about 1 m size or less. Bulk hardness also shows an increase as the titanium amount was increased (Fig. 18). For the iron with no titanium additions hardness was 48 HRC and increased to 58 HRC for the alloy with 2%Ti. It seems that the dispersion strengthening of matrix strongly influenced the increase in bulk hardness despite the decrease in the eutectic carbide volume fraction. In addition, the finer structure obtained by titanium additions may have also contributed to increase hardness. Results on Ti-added high-chromium chill-cast irons using wedgeshaped copper moulds [14] are in agreement with the results reported here. On the other hand, fracture toughness did not show a clear trend as the titanium √ amount increased in the iron (Fig. 18). It remained at 18 MPa m as a mean value for all titanium additions. An increase in bulk hardness could imply a decrease in fracture toughness; however, it is possible that matrix
Fig. 13. Agglomeration of TiC particles in some areas of the austenitic matrix in some titanium added high-chromium white cast irons: (a) 1.68%Ti, and (b) 2.02%Ti. Most particles are 1 m size maximum. SEM.
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Fig. 14. TEM dark field micrographs showing agglomeration of TiC (white) in the austenitic matrix (dark) with the zone axis [0 1¯ 1] oriented perpendicular to the surface as shown by the diffraction pattern (DP). Also some rings of randomly orientated TiC can be observed in the DP which correspond to the indexed planes.
Fig. 15. SEM micrographs showing coarse (3–5 m) TiC attached to the eutectic M7 C3 carbides. (a) 1.68%Ti alloy, and (b) 2.02%Ti.
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strengthening by particle dispersion has been compensated with the decrease in volume fraction of the brittle eutectic carbide and the structure refinement. Under these conditions fracture toughness did not show a change. On this basis, an improvement in wear resistance under sliding conditions is expected for this alloy.
4. Conclusions
Fig. 16. Eutectic carbide volume fraction as a function of the titanium content in the alloy.
• Titanium additions refined the structure of the alloy since the early TiC precipitation served as nuclei for proeutectic austenite. • A small reduction of the volume of eutectic M7 C3 carbide was observed due to the decrease of carbon after TiC was formed. TiC volume fraction was about 3% for the iron with 2%Ti. Eutectic carbide volume fraction was reduced from 0.25 for the base material to 0.21 for the iron with 2%Ti. • Titanium additions caused an increase in matrix microhardness due to the formation of hard small TiC particles within the matrix. • Bulk hardness was also increased mainly by the reinforcement of matrix by TiC, but the structure refinement may have also contributed. • Fracture toughness was not considerably affected as the titanium amount increased.
Acknowledgements
Fig. 17. Microhardness of matrix as a function of the titanium content in the alloys.
This research was supported by the Universidad Michoacana de San Nicolas de Hidalgo, Mexico. We are grateful for the financial support. One of the authors, R. Correa, acknowledges the National Council of Science and Technology from Mexico (CONACyT) for the supplied scholarship during the carrying out of this project.
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Fig. 18. Bulk hardness and fracture toughness as a function of the titanium content for the different alloys.
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Effect of titanium on the as-cast microstructure of a 16%chromium white iron A. Bedolla-Jacuinde ∗ , R. Correa, J.G. Quezada, C. Maldonado Instituto de Investigaciones Metal´urgicas, Edificio “U” Ciudad Universitaria, Universidad Michoacana de San Nicol´as de Hidalgo, Morelia Michoac´an, C.P. 58000, Mexico Received 17 August 2004; received in revised form 15 March 2005; accepted 29 March 2005
Abstract This research work studies the effect of systematic titanium additions (up to 2 wt.%) to a 16%Cr, 2.5%C white cast iron. The study was undertaken in six laboratory made alloys with different titanium amounts. Alloys were melted in an open induction furnace by using high purity raw materials. Such additions caused small hard titanium carbide particles to precipitate within the proeutectic austenite therefore promoting a strengthening of matrix; such particles also contributed to increase bulk hardness of the overall alloy. A structure refinement, as measured by the secondary dendrite arm spacing, was also observed as the titanium amount was increased. Titanium carbide precipitation caused a small decrease in the eutectic carbide volume fraction. The fracture toughness remained constant since the strengthening of matrix was compensated with a decrease in the volume of eutectic carbides. According to this study, titanium can be used as an alloying element to increase the hardness and perhaps wear resistance without affecting fracture toughness in high-chromium cast alloys. The results are discussed in terms of the precipitation nature of such small hard titanium particles. © 2005 Elsevier B.V. All rights reserved. Keywords: Cast iron; Titanium; Microstructure; High-chromium
1. Introduction High-chromium white irons can be described as in situ composites with large, hard eutectic and/or proeutectic M7 C3 carbides in a softer iron matrix (commonly austenitic in the as-cast conditions and mostly martensitic after a destabilization heat treatment). These microstructure combinations make high-chromium white irons very suitable for applications where abrasion resistance is the main requirement. Their exceptional abrasive and erosive wear resistance results primarily from their high volume fraction of hard carbides, although the toughness of the matrix also contributes to the wear resistance [1]. Attempts to improve the wear resistance in highchromium irons have led researchers to try different alloying elements such as vanadium, tungsten, titanium and niobium [2–14]. Even the addition of carbides such as TiC [15] has ∗
Corresponding author. E-mail address: [email protected] (A. Bedolla-Jacuinde).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.072
been undertaken to obtain composites in these irons. The aim of these additions is usually to achieve some modification of the eutectic carbide structure by obtaining harder carbides, though they may improve the hardenability of the matrix since these carbide forming elements may partition also to the matrix. Laird [16] and Ono et al. [17] have widely studied the partition coefficient or segregation ratio for several elements in high-chromium irons. It has been recognised that a possible strategy for improving the toughness of white iron alloys as well as the wear resistance under sliding conditions involves the refinement of the eutectic carbide structure by producing finer, more globular carbides [18–20]. Rapid cooling [21,22] or lower superheat [16] have been used as a means of increasing nucleation and hindering carbide growth, producing finer carbides. Conversely, slow cooling of castings results in larger dendrite arm spacing, reducing the number of sites for nucleation of the eutectic and hence, larger, coarser, eutectic forms [23,24]. Alloying additions of some elements that segregate to the matrix/carbide interphase, such as rare earths, silicon,
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boron, etc., have also been used to modify the eutectic carbide structure [25–34]. Generally, these attempts to modify the eutectic carbide structure through alloying elements have had limited success. Wear resistance may be increased by increasing the eutectic carbide volume fraction in the alloy but at the cost of further decreasing fracture toughness. Researchers, therefore, have the challenge of improving wear resistance without reducing fracture toughness. The high amount of eutectic carbides play a critical role in the fracture behavior of white iron alloys, and the toughness has been seen to deteriorate as the proportion of carbides in the structure increases [35]. Since a high proportion of carbides in the structure is required to maximise the wear resistance, many researchers have sought to improve the toughness through modifications to the carbide structure [26,32]. The carbide attributes which have been manipulated in an attempt to improve toughness include volume fraction [35], carbide size [36,37], and inter-carbide spacing, and carbide morphology and shape [36]. These studies have found, with varying degrees of success, that the toughness of high-chromium cast irons may be improved by decreasing the proportion of carbides in the structure, decreasing the carbide size, increasing the inter-carbide spacing, and globulizating the carbide shape. From this study it is expected to improve wear resistance in a 16%Cr white cast iron by adding titanium as an alloying element. Several carbide forming elements have been used for this purpose, as mentioned above, but the use of titanium has been very limited due to the difficulties found during ironmaking. Titanium is a very oxidation prone element and special conditions for alloying are required when using an open induction furnace (not controlled atmosphere). The hypotheses for using titanium to improve wear resistance without decreasing fracture toughness are: • Titanium is a strong carbide forming element; since TiC has a high formation temperature, it will be the first phase precipitating during solidification. Therefore, a proeutectic matrix reinforced with hard small TiC particles is expected at the end of the solidification process. The reinforced matrix, in turn, is expected to improve wear performance. • It is also expected that TiC formation refines the structure since such particles may act as nuclei for the austenite phase. • As a result of this, a fine microstructure with isolated hard small TiC within the matrix would be expected which improves wear performance.
2. Experimental procedure The alloys used in the present study were made in a laboratory induction furnace by using high purity raw materials. Thirty kilograms of the alloy were melted and poured at 1450 ◦ C into 5 kg capacity SiC crucibles where the titanium
additions were undertaken. Therefore, six crucibles with different additions of titanium were obtained which were then poured into sand moulds to obtain square 2.5 cm × 2.5 cm cross section cast bars. The amount of titanium was zero for the first ingot, 0.1% for the second, 0.7% for the third, 1.5% for the fourth, 1.7% for the fifth, and 2% for the sixth. Chemical analysis was undertaken by spectrometry from chill samples obtained during casting each ingot, and the results are shown in the next section. Samples for metallography, hardness, wear, and fracture toughness tests were cut from the cast bars. Sample cutting was carried out in a Discotom by using an abrasive wheel of alumina; cutting was done as slow as possible in order to avoid excessive overheating that may cause cracking in the samples. Additionally, copious amounts of water as coolant were used. Material’s characterization was undertaken by light microscopy, scanning electron microscopy (SEM), transmission electron microscopy (TEM), microanalysis by energy-dispersive spectrometry (EDS), X-ray diffraction (XRD) and PC based image analysis. Samples for metallography were prepared in the traditional way; roughed by using abrasive paper and then polished on nylon cloths by diamond paste. Once polished, the specimens were etched with one of the two different reagents: Villela’s reagent (5 ml HCl, 1 g picric acid in 100 ml ethanol) for 30 s to reveal the microstructure; or in a solution of 50 ml FeCl3 plus 20 ml HCl in 930 ml ethanol for about 3 h for deep etching. The later reagent removes part of the matrix without affecting the carbide phase providing a very high contrast between matrix and carbides when the structure is observed under the light microscope. Carbide volume fraction and carbide size were measured by image analysis using digital images taken on the light microscope at about 250×. For this purpose, samples were deep-etched to get high contrast. Twenty digital micrographs were processed using the software Sigmascan V5. Further characterization was undertaken by XRD in a SIEMENS 5000 diffractometer using Co K␣ radiation in a 2θ range of 30–130◦ . A JEOL 6400 SEM was also used for imaging and microanalysis. In addition, some thin foils for transmission electron microscopy were prepared from the iron with higher titanium amount. Thin sections were cut in discs of 3 mm and milled to 80 m, then dimpled to 15 m and finally ion milled to perforation. Samples were analyzed in a Philips Tecnai F20 TEM. Microhardness of the matrix in the different irons was measured on metallographic samples lightly etched with Vilella’s reagent. About 15 measurements for each sample were undertaken using a microhardness tester LECO M-40 with Vickers diamond indenter applying a load of 100 g for 15 s. Hardness, on the other hand, was measured on the Rockwell C scale for the bulk material. Between 10 and 15 measurements were undertaken on as-polished samples using a Rockwell hardness tester with 150 kg load and a diamond pyramid indenter typical for this scale.
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Table 1 Chemical composition of the high-chromium cast irons used for the analysis Element (wt.%)
Sample 1
Sample 2
Sample 3
Sample 4
Sample 5
Sample 6
C Cr Mo Ni Mn Si Ti
2.40 15.92 2.91 2.52 1.77 1.41 0.0
2.46 16.02 2.88 2.62 1.78 1.48 0. 11
2.48 16.02 2.92 2.49 1.90 1.4 0.72
2.45 15.82 2.95 2.53 1.86 1.45 1.53
2.61 16.12 2.90 2.55 1.84 1.45 1.68
2.58 15.90 2.89 2.60 1.77 1.45 2.02
Fracture toughness testing was carried out by three point bending in a Hounsfield universal testing machine. The samples whose dimensions were 2.5 cm × 2.5 cm × 15 cm, were U-notched and tested at a speed of 0.5 mm/min according to the ASTM E 399-90 standard [38] but without pre-cracking. The reported KIC values are the mean of three tests.
3. Results and discussion Table 1 shows the chemical composition for the six white irons used for the study. Note the similar amounts of carbon, chromium, nickel, molybdenum, manganese and silicon obtained for the alloys; the only varied element was titanium from 0 to 2.02%. Since the pouring temperature and cooling rate were the same for all the irons, variations in microstructure as well as in mechanical properties can be attributed to the effect of different titanium amounts. 3.1. Microstructure Fig. 1 shows an optical micrograph where the general microstructure of the iron with no titanium additions can be observed. The microstructure consists of hard eutectic carbides in a soft austenitic matrix. The solidification path fol-
lowed by this kind of alloys has been clearly summarized by Tabrett et al. [39]; being a hypoeutectic alloy, solidification starts with the formation of austenite dendrites which grow while the temperature decreases and the remaining liquid reaches the eutectic composition; then the eutectic reaction takes place and the coupled austenite/carbide eutectic is developed. Austenite remains as a metastable phase at room temperature due to the high amount of alloying elements, particularly molybdenum and nickel that increase hardenability and carbon that lowers the martensite start temperature MS at temperatures below room temperature [1]. XRD studies on this iron also show the presence of small amounts of martensite and the molybdenum rich M2 C carbide in addition to austenite and the eutectic M7 C3 carbide, Fig. 2. The presence of M2 C has been reported by several authors [1,16,17,25,40] in these alloys. According to them, most of molybdenum in the alloy partitions to the eutectic M7 C3 carbide and some dissolves in austenite (the Mo dissolved in austenite is the one that really helps to improve hardenability in the alloy), but a small part of Mo also forms a type M2 C carbide due to the presence of high amounts of carbon in the iron. Such a carbide forms as a eutectic and is the last part to solidify. It is commonly found joined together with the M7 C3 but it can be clearly distinguished from it from the lamellar shape of the former. Fig. 3 shows an SEM micrograph indicating the presence of the M2 C carbide and Fig. 4 shows detail of such a carbide in a deep-etched sample along with its composition, which clearly indicates is a molybdenum-rich carbide. The presence of martensite in these alloys has also been widely reported elsewhere [5,14,16,25]. During eutectic solidification the M7 C3 carbide which grows along the austenite, absorbs carbon from its surroundings and a narrow area at the austenite/carbide interphase becomes impoverished in carbon. The lack of carbon in these zones of austenite increases the MS temperature which allows these areas of austenite to transform to martensite during subsequent cooling down. Therefore it is common to find eutectic M7 C3 carbides surrounded by laths of martensite. Fig. 5 shows an SEM micrograph to confirm this phenomenon. 3.2. Effect of titanium on the microstructure
Fig. 1. Light micrograph showing the general microstructure of the iron without titanium additions. The structure is basically composed of eutectic M7 C3 carbides in an austenitic matrix. Vilella’s reagent for 30 s.
The microstructure of the irons with titanium additions is similar to the base iron, the only difference is the presence of titanium carbide (or carbonitride) mostly within the
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Fig. 2. XRD traces for the iron with no titanium additions showing the microstructure to consist of austenite, M7 C3 and small amounts of martensite and M2 C.
Fig. 3. SEM micrograph showing the presence of lamellar M2 C molybdenum-rich carbide.
Fig. 5. SEM micrograph of a high-chromium cast iron showing the presence of martensite laths at the carbide/matrix interface due to carbon depletion at these areas.
Fig. 4. SEM micrograph of a deep-etched sample showing detail of the M2 C carbide and the EDS analysis evidencing it is a molybdenum-rich carbide, the major element that determines the stoichiometry.
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Fig. 6. Sequence of SEM micrographs showing the increase in the amount of TiC (arrowed) within the austenitic matrix as the titanium content increased in the alloy.
proeutectic austenite matrix. Fig. 6 shows the microstructural evolution of the irons as the amount of titanium increases in the alloys. Note the increase in the number of small titanium carbides in the proeutectic austenitic matrix. XRD traces on the 2%Ti iron shows a couple of peaks indicating the presence of TiC (Fig. 7). Furthermore, chemical composition of these particles was clarified by EDS in the SEM (Fig. 8) and results confirm the particles to be TiC. In addition of titanium and carbon, some other elements such as molybdenum, chromium and iron are present in the particle, according to
the EDS. A more detailed characterization of these particles was undertaken by TEM as shown from Fig. 9. That figure shows a bright field micrograph where a hexagonal-shaped TiC particle can be observed, which was oriented to a zone axis [0 1 1] perpendicular to the plane; previous TEM calibration indicated good correlation between the measured distances from the transmitted beam to the indexed spots and the ratios obtained from databases. In addition, EDS analysis by a fine electron beam indicates only the presence of carbon and titanium in the particle. This later observation confirms
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Fig. 7. XRD traces for the iron with 2.02% titanium additions showing the microstructure to consist in austenite, M7 C3 , small amounts of martensite, M2 C and TiC.
that the presence of any other element in the EDS analysis undertaken by SEM might be due to the matrix interaction with the beam. Fig. 10 shows a bright field TEM micrograph where a TiC particle appears white and the dark austenitic ¯ as shown by the matrix is oriented with the zone axis [1 1¯ 2] diffraction pattern. In addition, EDS on the matrix indicates the presence of elements such as chromium, molybdenum, nickel, manganese, silicon, carbon and titanium. Cr, Mo, Fe, C and Mn also may be found in the carbide phase since they partition to both carbide and matrix phases [14,17,23,39,41]. From the present study, the specific temperature at which these fine TiC particles (1 m size or less) formed was not determined. However, it may be deduced that these fine particles formed at some temperature higher than that of austenite dendrite formation (when the alloy is still liquid) and that these small carbides were entrapped by the growing austenite dendrites. Since these TiC particles are the first to solidify, it is likely they serve as nuclei for proeutectic austenite. Actually, a refinement of the secondary dendrite arm spacing from 25
to 18 m was observed when titanium was increased from 0 to 2.02% (Fig. 11). Such a refinement has also been observed and reported in a previous work [14] and is attributed to the precipitation of TiC particles when the alloy is still liquid. If these particles actually function as nuclei for austenite, there must be a crystallographic relationship between some planes of TiC and austenite; however, such a relationship was not studied in the present work. Fig. 12 shows an SEM micrograph of the 2%Ti iron where a uniform particle distribution within the matrix is observed. However, such a distribution was not observed all around the sample. Some areas where agglomeration of TiC particles occurred were also observed in the samples with high Ti additions (see Fig. 13). Fig. 14 shows a mosaic of dark field micrographs of agglomerated TiC particles where austenite appears dark and TiCs are white (the image was generated by the diffracted [2 0 0] spot of TiC indicated in the diffraction pattern). Also a diffraction pattern of the area in the mosaic was undertaken, which shows spots of austenite and
Fig. 8. SEM micrograph showing two TiC particles in austenite and the EDS showing the chemical composition of the arrowed particle. 2%Ti iron sample, Vilella reagent, 45 s.
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Fig. 9. Bright field TEM micrograph showing a hexagonal-shaped TiC particle (a) with its corresponding selected area diffraction pattern (b) and chemical composition as undertaken by EDS (c). Sample of the 2%Ti iron.
rings formed for the several TiC particles in the area. Such rings are not continuous due to the large size of particle and perhaps some degree of preferred orientation. This last observation reinforces the theory that there must be some crystallographic relationship between austenite and TiC, so that austenite could precipitate on the formed TiC. However, not all the particles whose size is about 1 m were well dispersed in the matrix; Fig. 15 shows SEM micrographs where coarser (up to 4 m size) TiC appeared attached to the eutectic carbides at the M7 C3 /austenite interface. Such a phenomenon may be explained as follows: a large number of TiC particles precipitate at high temperatures and most of them serve, perhaps, as nuclei for the subsequent austenite
formation. Once primary austenite dendrites have formed, they will grow rejecting some alloying elements into the liquid and collide with some other TiC particles. Since titanium has a low partitioning coefficient to austenite [42], a titanium rich layer is expected at the periphery of the austenite arms. That increase in solute probably induces a constitutional supercooling which facilitates coarsening of particles. The carbon needed for the carbide formation is easily removed from the melt on account of its high diffusivity. Thus, whenever the free carbide particles collide with the interface, they are pushed away by the dense layer. Sometimes, these particles are caught by the faster moving interface and appear at the centre of the dendrites with smaller sizes; they may have been
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Fig. 10. Bright field TEM micrograph showing a hexagonal-shaped TiC particle in matrix (a), (b) selected area diffraction pattern of matrix showing to be austenite, and (c) the chemical composition of matrix by EDS showing the presence of dissolved elements such as Cr, Ni, Mo, C, Mn, Si and even Ti. Sample of the 2%Ti iron.
entrapped at the early stages of their coarsening or they nucleated austenite. Kesri and Durand-Charre [7] have reported similar observation for the NbC formation during solidification of Fe–Cr–C alloys. Fig. 16 shows variations in the eutectic carbide volume fraction as the amount of titanium increased in the irons. It is well known that in iron–carbon alloys, particularly in
cast irons, high solidification rates favor the formation of the metastable carbide/austenite eutectic instead of the stable austenite/graphite eutectic independently of the presence of other graphitizing or carbide forming elements. For the present study, all the alloys were solidified at the same rate and the differences observed are attributed to the effect of titanium. Such differences are due to the precipitation of TiC
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prior to the eutectic solidification. The formation of these carbides consumes part of the carbon present in the alloy so that during eutectic solidification the volume of M7 C3 is diminished. These results are in agreement with some works where titanium and/or niobium have been used as alloying elements in high-chromium irons [3,7,11,14,41]. 3.3. Effect of titanium on the mechanical properties
Fig. 11. Secondary dendrite arm spacing as a function of the titanium content in the alloy.
Fig. 12. SEM micrograph of the 2%Ti white iron showing the well distributed small TiC particles in the austenitic matrix.
Fig. 17 shows the increase of matrix hardness as the amount of titanium increases in the alloy. This logical phenomenon is achieved due to the reinforcement of austenite by a large number of hard small well dispersed TiC particles. It is also possible that titanium dissolved in matrix for the irons with higher Ti amounts (see Fig. 10), may have contributed to the matrix strengthening by solid solution; however, such a contribution would be negligible compared to the dispersion strengthening. For this case, a considerably increase in hardness from 330 to 500 Vickers100 was observed. The maximum volume of TiC was ∼3% for the iron with 2%Ti with particles about 1 m size or less. Bulk hardness also shows an increase as the titanium amount was increased (Fig. 18). For the iron with no titanium additions hardness was 48 HRC and increased to 58 HRC for the alloy with 2%Ti. It seems that the dispersion strengthening of matrix strongly influenced the increase in bulk hardness despite the decrease in the eutectic carbide volume fraction. In addition, the finer structure obtained by titanium additions may have also contributed to increase hardness. Results on Ti-added high-chromium chill-cast irons using wedgeshaped copper moulds [14] are in agreement with the results reported here. On the other hand, fracture toughness did not show a clear trend as the titanium √ amount increased in the iron (Fig. 18). It remained at 18 MPa m as a mean value for all titanium additions. An increase in bulk hardness could imply a decrease in fracture toughness; however, it is possible that matrix
Fig. 13. Agglomeration of TiC particles in some areas of the austenitic matrix in some titanium added high-chromium white cast irons: (a) 1.68%Ti, and (b) 2.02%Ti. Most particles are 1 m size maximum. SEM.
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Fig. 14. TEM dark field micrographs showing agglomeration of TiC (white) in the austenitic matrix (dark) with the zone axis [0 1¯ 1] oriented perpendicular to the surface as shown by the diffraction pattern (DP). Also some rings of randomly orientated TiC can be observed in the DP which correspond to the indexed planes.
Fig. 15. SEM micrographs showing coarse (3–5 m) TiC attached to the eutectic M7 C3 carbides. (a) 1.68%Ti alloy, and (b) 2.02%Ti.
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strengthening by particle dispersion has been compensated with the decrease in volume fraction of the brittle eutectic carbide and the structure refinement. Under these conditions fracture toughness did not show a change. On this basis, an improvement in wear resistance under sliding conditions is expected for this alloy.
4. Conclusions
Fig. 16. Eutectic carbide volume fraction as a function of the titanium content in the alloy.
• Titanium additions refined the structure of the alloy since the early TiC precipitation served as nuclei for proeutectic austenite. • A small reduction of the volume of eutectic M7 C3 carbide was observed due to the decrease of carbon after TiC was formed. TiC volume fraction was about 3% for the iron with 2%Ti. Eutectic carbide volume fraction was reduced from 0.25 for the base material to 0.21 for the iron with 2%Ti. • Titanium additions caused an increase in matrix microhardness due to the formation of hard small TiC particles within the matrix. • Bulk hardness was also increased mainly by the reinforcement of matrix by TiC, but the structure refinement may have also contributed. • Fracture toughness was not considerably affected as the titanium amount increased.
Acknowledgements
Fig. 17. Microhardness of matrix as a function of the titanium content in the alloys.
This research was supported by the Universidad Michoacana de San Nicolas de Hidalgo, Mexico. We are grateful for the financial support. One of the authors, R. Correa, acknowledges the National Council of Science and Technology from Mexico (CONACyT) for the supplied scholarship during the carrying out of this project.
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