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Mechanochemical reaction between ilmenite (FeTiO3) and aluminium
Journal of Alloys and Compounds 270 (1998) 228–236
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Mechanochemical reaction between ilmenite (FeTiO 3 ) and aluminium N.J. Welham* Petrochemistry and Experimental Petrology, Research School of Earth Sciences, and Department of Electronic Materials Engineering, Research School of Physical Sciences and Engineering, Australian National University, Canberra ACT 0200, Australia Received 2 December 1997
Abstract A natural ilmenite (FeTiO 3 ) and aluminium powder have been mechanically milled together for 100 h in a laboratory ball mill. The as-milled powder and an unmilled powder of identical composition were annealed at up to 12008C and examined by X-ray diffraction and differential thermal analysis (DTA). The unmilled sample showed aluminium melted prior to an exothermic reaction starting at |8508C. The milled powder showed no thermal activity, other than a reversible phase transition at 1067648C, indicating that reaction occurred within the mill. The products of both powders were the same, TiAl 3 , Fe 4 Al 13 and Al 2 O 3 , although in the milled powder these phases were nanocrystalline until annealing caused crystallite growth. The thermal reaction seemed to occur in two stages, formation of TiAl 3 , Al 2 O 3 and elemental iron followed by a slower, diffusion controlled reaction between the elemental iron and residual aluminium to form Fe 4 Al 13 . The reaction during milling was attributed to increased intermixing between the ilmenite and aluminium causing a change in the rate determining step from solid-state diffusion to another, unknown, controlling mechanism. 1998 Elsevier Science S.A. Keywords: Aluminium metal-oxide systems; Mechanical milling; Solid state diffusion
1. Introduction There has been increasing interest in the properties and uses of intermetallic compounds over the past few years, primarily due to their superior properties when compared with those of the constituent metals. It is well known that dispersion of a hard material e.g. SiC, Al 2 O 3 within the metal matrix can give enhanced mechanical properties over that of the pure intermetallic compound. Particles are often used as the hard phase, although fibres are more common, giving strength improvements of 30% in some cases [1], the particles can be formed by in situ thermal reaction [2] or added during a further processing stage such as infiltration [3] or milling [4,5]. Several aluminium–metal oxide systems have been milled, leading to reduction within the mill to form an alumina dispersion in the metal [6–8]. A previous paper [9] has indicated that the reaction between aluminium and TiO 2 to form alumina dispersed in TiAl 3 has its thermal onset temperature lowered by 5008C by milling the powders prior to heating. Additionally, the reaction became potentially self-propagating after milling, whereas it was net endothermic without milling. There may also be advantages in using mixtures of *Tel.: 0061 2 6249 0520; fax: 0061 2 6249 0511; e-mail: [email protected] 0925-8388 / 98 / $19.00 1998 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 98 )00123-6
intermetallic compounds [1,10,11], although the specific gains will vary with the combination used. This paper examines the milling of a mixture of naturally occurring ilmenite and aluminium to determine whether a composite of alumina dispersed in a mixture of iron and titanium aluminides can be formed at low temperature.
2. Experimental The ilmenite concentrate was the same as that detailed previously [12] and contained ,5% impurities. The asreceived particle size was 90% 210690 mm, but the powder was milled in vacuum for 25 h prior to use to reduce the particle size to .95% ,2 mm [13]. The aluminium powder was nominally .99% Al, although some surface oxide was evident, the particle size was ,20 mm. Feed powders were prepared in accordance with the stoichiometry of reaction (1). This stoichiometry was chosen on the basis of thermodynamic predictions, using free energy minimisation, which showed 7:1 to be the ratio beyond which additional aluminium remained unreacted. However, this prediction was made using species for which thermodynamic data was available in the database [14] with the addition of data at 298.15 K for the iron
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aluminides FeAl, FeAl 2 and FeAl 3 [15]. Since no data was found for the phases Fe 3 Al, Fe 2 Al 5 and Fe 4 Al 13 this reaction can only be considered a guide. FeTiO 3 1 7Al ⇒ FeAl 2 1 TiAl 3 1 2Al 2 O 3
(1)
Duplicate samples of each mixture (with a nominal 10 mol % excess of aluminium to account for surface oxide) were sealed in separate laboratory-scale ball mills. Five 10 (25.4 mm) diameter stainless steel balls were added to one mill giving a ball:powder mass ratio of 47:1. The mills were evacuated to ¯10 Pa and rotated at 165 rpm for 100 h, using magnets to control ball motion within the mill [16]. After milling, there was no change in gas pressure, indicating that gas was neither evolved nor leaked into the system. The sample derived from the mill which contained balls was considered to be milled, that from the mill with no balls was unmilled. Differential thermal analysis (DTA) was performed on the milled and unmilled powders. Samples of 3062 mg were loaded into an alumina crucible and heated up in an inert argon atmosphere to 12008C and cooled to 5008C, both at 208C min 21 using a Shimadzu DTA-50 instrument. Isothermal experiments were also performed on the unmilled powder to determine the rate controlling step by modelling the DTA data [17–20]. For the optimised model at each temperature, values of the rate constant were plotted against reciprocal temperature and activation energy of reaction determined. Thermogravimetric analysis (TGA) was performed under identical conditions in a Shimadzu TGA-50. Samples of powder were also isothermally annealed under flowing argon for 1 h. The products were analysed by X-ray diffraction (XRD) ˚ using monochromatised Co Ka radiation ( l 51.78896 A) using a count time of 2 s per 0.018 step over the range 10–1108. A narrower range of 20–708 is presented in the figures to aid clarity, but calculations and identification were made using the complete range of 2u measured. The crystallite size was estimated by applying the Scherrer equation [21] to all of the available identified peaks for each phase. The error reported in the crystallite size is 2s for the available data. Scanning electron microscopy was used to determine morphology and the distribution of the iron, aluminium and titanium within selected powders.
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shown in Fig. 1(a), no thermal reactions were detected below this temperature for any powder. During the heating of the unmilled powder there is a large endotherm around 6608C, this is due to the melting of the aluminium in the powder. The energy associated with this melting was equivalent to 0.36 kJ / g Al, which is 10% lower than the accepted value of the enthalpy of fusion for Al (0.397 kJ / g Al) [22]. This would indicate that |10% of the mass of the aluminium powder was present as oxide, indicating that the apparent 10% molar excess of aluminium used should give almost the exact stoichiometry for reaction (1). After the aluminium had melted, there was an exotherm which started at 8658C and was completed by 10908C. The energy associated with this exotherm was 0.32 J / g of powder (0.55 kJ / g Al) which is considerably less than the 1.0 kJ / g necessary [14] to heat the mixture of ilmenite and aluminium up to the ignition point at 8658C indicating that the reaction cannot be self-propagating. A slower heating rate of 58C min 21 showed that the peak was due to a single reaction. On cooling the system, the only event was the freezing
3. Results
3.1. Unmilled powder TGA of unmilled powder showed that there were no mass changes (,0.5%) when heated up to 12008C, confirming the entirely solid state nature of the reactions occurring, as expected from reaction (1). DTA traces of the unmilled powder for T .5008C is
Fig. 1. DTA traces, normalised to equal mass, for powders heated to 12008C at 208C min 21 and cooled at 208C min 21 to 5008C under argon. The dashed lines represent the cooling portion of the cycle. (a) unmilled, first cycle; (b) unmilled, second cycle and (c) milled for 100 h.
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of residual aluminium, the energy evolved indicated that |24% of the aluminium had not reacted during the heating excursion. Trace (b) in Fig. 1 shows the results of heating the residue of trace (a) up to 12008C again. Clearly, the only major thermal event was the melting and freezing of the unreacted aluminium. The energy associated with the freezing was only 85% that of the melting suggesting that aluminium was being consumed by a reaction above the melting point, this was confirmed by subsequent heating / cooling cycles which showed a slow decrease in the peaks associated with elemental aluminium. The absence of a clear thermal event involving aluminium was most probably due to it occurring slowly, giving a rise in the differential temperature which was similar to that of the background noise. This would suggest that reduction within the system was completed during the initial heating or, an exotherm associated with alumina formation would be expected due to its large enthalpy of formation from aluminium (21676 kJ mol 21 at 298 K). The slowness of this second reaction would explain the diminution of residual aluminium with the number of heating excursions up to 12008C. Indeed, DTA of a powder annealed at 12008C for 1 h showed a similar trace to that of Fig. 1(c), but with the addition of a small pair of peaks for melting / freezing of aluminium. Clearly, even 1 h at 12008C was insufficient for complete reaction to occur, although the annealed powder was becoming similar to that of the powder milled for 100 h. Isothermal DTA was performed on the unmilled powder at several temperatures within the exotherm. Modelling [17–20,23] of the exotherm showed that an Avrami– Erofeev model with exponent 0.67 was the best fit, as demonstrated by the plot inset in Fig. 2. The physical meaning of this model is that the reaction rate is controlled by the movement of the reaction interface in three dimensions [23]. The Arrhenius plot, Fig. 2, shows that the activation energy for the reaction was 358 kJ mol 21 , somewhat lower than the values of 517 and 483 kJ mol 21 found for the formation of TiAl 3 from elemental aluminium and titanium [24,25]. The latter of these values was obtained using a second order reaction model, which has been considered to have no physical meaning in solid-state reactions [23,26] and, hence is open to uncertainty. The onset temperature for the reduction step (8808C) was somewhat greater than that for synthesis from the elements (7308C) [24]. The clear implication is that the controlling step of reaction (1) was not formation from the elements, as would be expected since thermal reduction of TiO 2 by excess Al did not show any evidence for the formation of elemental titanium [9]. Examination of the XRD trace of the unmilled powder, shown in Fig. 3(a), indicates that aluminium and ilmenite are the only phases present. The aluminium shows narrow peaks indicating a large crystallite size (7366 nm), whereas the ilmenite shows somewhat broader, less intense peaks typical of a phase of small crystallite size (17610
Fig. 2. Arrhenius plot of the rate constant for the data from unmilled powder isothermally heated in the DTA. The rate constant was derived from the mathematical model of da / dt versus t based on the Avrami– Erofeev equation with exponent 0.67. A sample plot of the model, for 9508C, is shown inset, with the hollow squares the data points and the line the model, the data points shown are 20% of those used in the fitting to aid clarity.
nm). This refinement of ilmenite crystallite size was due to the 25 h premilling prior to mixing. A sample of the unmilled powder was annealed at 6008C for 1 h, the XRD trace of this powder is shown in Fig. 3(b). There were no new peaks present, indicating that no apparent reaction had occurred. The only change was a narrowing of the peaks due to thermally induced crystallite growth, giving 2868 nm and 8869 nm for ilmenite and aluminium respectively. A broad peak around 318 was also evident, this was probably due to rutile (TiO 2 ). The initial ilmenite powder contained a small amount of rutile [12] and its presence may be due to crystallite growth, which has been noted to occur below 6008C [9]. Rutile is also the initial reduction product of ilmenite [27–31] and its presence could also be due to reduction. The other product of ilmenite reduction is elemental iron, unfortunately, the main peaks for elemental iron are coangular with peaks of aluminium and cannot be deconvoluted, primarily due to the low intensity of the iron peaks that would be expected
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Fig. 3. XRD traces for powders: (a) unmilled; (b) unmilled, 1 h at 6008C; (c) unmilled, 1 h at 10008C; (d) 100 h milled; (e) 100 h milled, 1 h at 6008C; (f) 100 h milled, 1 h at 10008C; (g) 100 h milled, 12 h at 12008C; m-Al, ♦-Al 2 O 3 , j-TiAl 3 , d-Fe 4 Al 13 , n-FeTiO 3 , s-Fe and h-TiO 2 data for Fe 4 Al 13 is only available for 2u #55.038.
to be present. However, the powder annealed at 6008C proved to be magnetic, the unmilled powder having showed no signs of magnetism. Clearly, some reduction reaction had occurred, presumably forming elemental iron, although the extent was probably small as there were no traces of an exotherm which would be expected when Al 2 O 3 forms from Al. The reaction occurred below the melting point of aluminium and can only have occurred at the contact points between the large aluminium grains and the smaller ilmenite particles. Annealing above the melting point of aluminium but below the exotherm, also gave a magnetic powder, XRD of which showed weak peaks where the main peaks of alumina and TiAl 3 were expected in addition to those of aluminium and ilmenite. Clearly, the melting of aluminium causes an increase in contact area between the reactant phases leading to a larger extent of reaction which allows detection by XRD. Only after annealing above the exotherm, were major phase changes observed in the XRD trace, as in trace (c). Both TiAl 3 and Al 2 O 3 emerged, the main aluminium peak
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had weakened considerably and ilmenite was absent, thus confirming that reaction (1) had occurred to some extent. The powder was magnetic, implying that elemental iron was still present, but its XRD peaks concealed. The rutile present at 6008C had been consumed by reduction. There were several peaks which were probably due to Fe 4 Al 13 , the most intense of these appearing as shoulders on the aluminium and alumina peaks at 51 and 538. The continued presence of aluminium after 1 h of annealing above the exotherm is a clear indicator that the formation of one phase is slow. The evidence seems to be that TiAl 3 is readily formed, with clear peaks present after 1 h at 10008C, Fig. 3(c). However, no iron bearing phase can be confirmed as present, although peaks are probably present for Fe 4 Al 13 . This would imply that the final iron bearing phase was the one which formed slowly. The free energy of formation of TiAl 3 from its elements is 2170 kJ mol 21 whereas that for FeAl 3 is 279 kJ mol 21 (the closest phase to Fe 4 Al 13 for which any data was available), clearly the formation of TiAl 3 is more favourable. The reduction sequence for ilmenite passes through elemental iron and TiO 2 [27–31] and it may be expected that the aluminium would react with the elemental iron in preference to the TiO 2 . However, the free energy for the reduction of TiO 2 by Al to alumina and titanium trialuminide is 2301 kJ mol 21 TiO 2 —much more favourable than reaction of aluminium with iron to form iron aluminide. There was no XRD evidence for intermediate phases, even after quenching DTA runs at points on the reaction exotherm corresponding to |10, 25 and 50% aluminium consumption. The magnetism evident in the annealed powders would seem to indicate that elemental iron was formed during the reduction and remained in the powder after 1 h at 10008C. Thus, it would seem that the rate determining step is the reaction of aluminium with the titanium phase from ilmenite reduction. Previous work [9] has shown that the thermally induced reaction between a mixture of aluminium and rutile, prepared identically to that used here, occurs at 10008C. Thus, it appears that the titanium product of ilmenite reduction was not crystalline rutile, or the main reaction exotherm would have been observed at higher temperature. The titanium phase formed may have been nanocrystalline and thus, not readily detectable by XRD. The separation of iron and titanium at the atomic level in ilmenite could cause more intimate mixing between the titanium phase and aluminium than in physical mixtures, lowering the reaction temperature as supported by the onset temperature of 6008C for milled rutile and aluminium [9]. The presence of elemental iron may also act as a catalyst by forming intermediate phases which reduces the energetic requirements. It has been shown that Fe 2 TiO 5 forms by reaction between elemental iron and TiO 2 at temperatures above 6008C [32,33] and phases of this type may be more amenable to reduction than TiO 2 . Regardless of the nature of the initial rate controlling
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step, the temperature required for this step to occur is greater than that necessary for the subsequent reactions forming alumina and TiAl 3 which consequently occur extremely rapidly. The final iron phase, Fe 4 Al 13 , would seem to be the phase formed relatively slowly as it was only detected after annealing for 1 h at 10008C. Its formation is presumably by interdiffusion of elemental iron and aluminium which would be expected to be slow at low temperatures giving a phase of varying composition and low crystallinity, which would be difficult to detect by XRD. With increased annealing temperature, the quantity of Fe 4 Al 13 formed increased and crystallite growth occurred, rendering it more readily detectable by XRD. Indeed, after 12 h at 12008C the powder was no longer magnetic showing that elemental iron had been completely consumed by reaction, no peaks for aluminium were evident in the DTA. The slow diminution of the aluminium melting / freezing peaks in the DTA with thermal cycling would support the slow consumption of aluminium, the overlap of the XRD peaks for iron and aluminium concealing the concomitant consumption of elemental iron.
3.2. Milled powder The DTA trace for the milled powder is shown in Fig. 1(c), the complete absence of any aluminium was clear, with no evidence for a melting endotherm. This proves that reactions involving elemental aluminium were completed within the mill. The only thermal events observed were the endotherm and exotherm of a phase change around 10708C. It is interesting to note the difference in shape between the heating and cooling events. The heating shows a symmetrical peak, but on cooling there are two overlapping peaks. The energy given out during cooling was around 95% that taken in during heating, this ratio was constant within error for all samples. A slower heating rate (58C min 21 ) was used to determine whether the peaks were separate, this showed essentially the same result—a symmetrical peak during heating and an asymmetric cooling peak. A sample of milled powder annealed at 12008C for 1 h showed an identical DTA trace to that presented in (c). The DTA trace of a sample of powder heated up to 10508C, just below the onset of the peak during heating did not show any peaks on cooling, whereas a sample heated to 10708C, just above the onset showed an asymmetrical cooling peak. This result confirms that the asymmetry was associated with the phase change, however, the reason for this is uncertain. It is possible that there are two phases transitions at slightly different temperatures, however the absence of separate peaks during heating would seem to discount this. Separation of the two overlapping exotherms observed during cooling was made assuming both were symmetrical, the energy associated with the second peak was 30% that of the main freezing peak. The transition point of the phase was estimated as the
mean of the temperatures at the maxima of the melting and freezing peaks [17,18,20], the average of four separate experimental runs gave a value of 1067648C. The XRD trace for the milled powder, Fig. 3(d), shows a broad peak at 468 and a very broad peak between 48–558, the first of these peaks was assigned to TiAl 3 , the absence of any other significant peaks makes this tentative. The second peak was too broad for profile fitting to reveal any profitable information, although the main peaks of several possible phases (e.g. Al 2 O 3 , FeAl, FeAl 2 , FeAl 3 , Fe 4 Al 13 ) are around this area and a combination of some, or all, of these may be present. The estimated crystallite size for the peak at 468 was |6 nm, but this is subject to potentially gross error due to uncertainties in whether this peak was for a single phase or was due to multiple phases with overlapping peaks. There was no evidence for either aluminium or ilmenite, although this may simply be due to reduction of the crystallite size, or amorphisation during milling. After annealing the milled powder at 6008C the trace (e) was obtained, crystallite growth is clearly evident for the main peak and is almost certainly due to TiAl 3 . Smaller peaks are now present for alumina but other peaks cannot be attributed with any certainty, particularly around 50– 558 where a broad peak is evident. The absence of peaks for ilmenite and aluminium can now be definitely attributed to reaction since the unmilled powder (b) showed crystallite growth occurred for both of these phases at 6008C. The XRD trace for the 1 h of annealing at 10008C is shown in trace (f), in which the grain growth of several phases is evident. TiAl 3 and Al 2 O 3 both show several distinct peaks thus confirming their presence. The majority of the remaining peaks were due to Fe 4 Al 13 , the main peaks for this phase are around 51–538 and these were almost certainly contributing to the broad peak noted in (d) and (e). The confirmation of the presence of Fe 4 Al 13 in the milled sample annealed to 10008C, would indicate that it was probably also present in the unmilled sample (c). The somewhat weaker peaks for the unmilled powder are due to the low fraction of the phase present, due to solid-state diffusion limitation of its formation. Unlike the unmilled powders no magnetic material was present in the as-milled powder, nor in any of the annealed powders. Clearly, milling performed the same function as high temperature annealing of unmilled powders, giving complete reaction between ilmenite and aluminium directly after milling.
3.3. Electron microscopy The secondary electron images and elemental maps for unmilled and milled powders are shown in Fig. 4(a) and (b) respectively. The unmilled powder consisted of large rounded grains surrounded by smaller submicron grains. The elemental maps confirm that the large grains were aluminium and the small grains ilmenite, as expected from
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Fig. 4. Micrographs (scale bar is 20 mm) and their corresponding digitally collected elemental maps of titanium, iron and aluminium for powders: (a) unmilled; (b) 100 h milled; (c) unmilled, 1 h at 10008C; (d) 100 h milled, 1 h at 10008C.
the feed powder particle size. The two phases were essentially discrete, although the ilmenite did appear to form a layer on the surface of the aluminium grains, due to adhesion between the particles during tumbling in the mill. The milled powder showed somewhat different morphology, the smallest grains are not evident, giving a minimum particle size around 1 mm, although the majority are 5–15 mm. The grains were comparatively angular consisting of submicron fragments which appear to have been welded together, forming larger agglomerates. An equilibrium particle size has been shown to be a typical feature of extended milling [13,34] although different materials / mixtures attain different particle sizes. It has been noted that the presence of a metallic phase, either added initially or formed by reaction during milling, tends to give a larger particle size than in a metal free system [35–37] indicating that the metal acts as a binder for the submicron particles. Elemental maps show that the aluminium, titanium and
iron were present in all grains, showing that milling produces intimate intermixing beyond the resolution of the mapping (512 mapping points per 61 mm giving an apparent resolution of 0.12 mm, but probably closer to |0.5 mm due to the interaction volume of the beam). The reaction that took place during milling could be expected to lead to this type of intermixing, although intermixing on a submicron scale has been demonstrated previously in a system where reaction did not occur during milling [9]. After annealing, the unmilled powder showed no significant change in particle size, Fig. 4(c), with large grains and small grains still evident. The large grains appear to be more angular and faceted than those of the unheated powder (a), this is probably due to crystallisation of aluminium from the liquid phase during cooling from 10008C. Elemental mapping shows that aluminium was also present in some of the small grains, indicating that reaction had caused some breakage of the large grains. Titanium and iron were still dispersed throughout the
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powder, although several grains show separate areas for titanium and iron, indicating some segregation of iron and titanium had occurred. There are a few areas where aluminium seems to be associated with the titanium as opposed to evenly distributed between the iron and titanium. This may explain why the formation of Fe 4 Al 13 was slow, as segregated areas of aluminium and iron would increase the diffusional pathlength, compared with the well mixed areas. The micrographs and elemental maps for the milled powder annealed at 10008C for 1 h are shown in Fig. 4(d). There seems to have been an increase in the maximum particle size compared with the unannealed powders (b) and the particles appear to be more angular. The elemental maps showed the expected result, with no evidence of segregation of iron, titanium and aluminium.
4. Discussion It is clear that the same reaction occurred during mechanical milling as during thermal treatment of unmilled power of identical composition. This could imply that the instantaneous temperature during impact reached that necessary for reaction i.e. .8508C. However, modelling of the thermally induced reaction for the unmilled powder indicated that the activation energy was 358 kJ mol 21 , typical of solid-state diffusional control. Elemental mapping of the powders showed that milling caused intimate mixing of the phases, not evident in the unmilled powder. Clearly, milling must have decreased the diffusional pathlength, hence causing a change in the rate controlling step. However, the absence of thermal reaction in the milled powder does not allow the activation energy of reaction for the milled powder to be determined. In both the milled and unmilled powders the final products after annealing were the same, TiAl 3 , Fe 4 Al 13 and alumina, although the extent of reaction was greater for the milled powder than the unmilled with residual aluminium in the unmilled powder. The thermal reaction seemed to occur in two main stages, the formation of alumina, TiAl 3 and elemental iron, followed by a slow diffusion controlled reaction between residual aluminium and iron to form Fe 4 Al 13 . A third reaction between elemental iron and TiAl 3 was also evident, but only after extensive annealing at high temperature. For the milled powder, the main reactions occurred within the mill, there being no residual aluminium detected in the DTA during initial heating. For milled samples pretreated in different ways, the only thermal process detected was a reversible phase transition at around 10678C, therefore elemental aluminium was definitely absent. Examination of the phase diagrams of the binary systems for the non-oxides that were present [38–41] showed no phases had melting points close to this value. A ternary diagram for the Fe–Al–Ti system [42] did not show any evidence for a three
component phase with a phase transition, although at least two further phases within this system (Ti 0.75 Fe 0.25 Al 3 and Fe 4 Ti 0.93 Al 12.07 ) have since been crystallographically described [43,44], but were not present. Examination by XRD provided no clues as to the identity of this phase, with only phases of higher melting point detected. A sample of unmilled powder, annealed for 1 h at 12008C showed similar DTA behaviour to that of the as-milled powder, with a transition point of |10708C evident. The presence of Fe 4 Al 11 would seem to be necessary for this energetic event to be detected. Phase diagrams for the iron–aluminium system [38,39] indicate a melting point around 11508C for a phase of this composition, but no phase transition has been reported. The overlapping peaks observed on cooling in the DTA may be due to separate phases of similar composition. During heating, the phases undergo transformation at a similar temperature that cannot be separated by DTA. Above the transition temperature, a diffusion reaction occurs which changes the relative stoichiometry of the phases slightly, but insufficient to change the shape of the DTA trace on cooling, one phase reversed its transition at a slightly higher temperature than the other giving rise to two peaks. A sample of milled powder was annealed at 12008C for 12 h, DTA on this sample showed a transition point at |10708C that did not show asymmetry on cooling, clearly, this long annealing time was sufficient for homogenisation to occur, giving a single uniform stoichiometry. The original stoichiometry used contains insufficient aluminium for complete formation of Fe 4 Al 13 , TiAl 3 and Al 2 O 3 . A ratio of 8.25 Al per ilmenite is required for complete reaction, whereas the amount used was |7 Al per ilmenite (accounting for 10% of oxide). Assuming that TiAl 3 and alumina form and all remaining Al forms Fe 4 Al 13 then 38.5% of the iron should be present as metal. This should be readily distinguishable from aluminium by the absence of a peak at |458 whilst the peaks at larger 2u should remain, however, XRD data for Fe 4 Al 13 is only available for 2u #55.038 and overlap may occur. A sample of powder which contained Fe 4 Al 13 did not show peaks at the larger 2u peaks of iron, thus confirming that the higher 2u peaks did not overlap and were diagnostic for iron. Iron was certainly present in the unmilled powder products as they were all magnetic, none of the milled powders were magnetic showing that, although there was excess iron in the system, it has been rendered nonmagnetic, presumably by reaction with aluminium to form an aluminide of uncertain composition. This iron aluminide could be the reason for the second peak during cooling of the DTA. The available thermodynamic data indicates that iron may react with the higher aluminides to form less aluminium rich aluminides, e.g. by reaction (2) or (3) which have respective free energies of 271 and 213 kJ mol 21 Fe at 258C. Similar calculations for reaction between titanium tri-aluminide and iron showed that reaction (4) was only just favourable at 258C with a free energy of 21
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kJ mol Fe, and hence would probably be a minor reaction compared with (2) and (3). The free energy at higher temperatures, where interdiffusion would be significant, could not be calculated due to the absence of entropy and heat capacity data for the iron aluminides in the data source used [15]. 2FeAl 3 1 Fe ⇒ 3FeAl 2
(2)
FeAl 2 1 Fe ⇒ 2FeAl
(3)
2Fe 1 TiAl 3 ⇒ TiAl 1 2FeAl
(4)
Milled and unmilled powders were annealed for 12 h at 12008C under argon to assess the possibility of reactions of these type occurring. The XRD traces of the resultant powders were essentially identical, with only a slight change in the relative intensities of the peaks. The milled powder trace, Fig. 3(g), showed only a weak TiAl 3 peak, the peaks for Al 2 O 3 and Fe 4 Al 13 remained approximately constant. Peaks at higher 2u which correspond to iron (and aluminium) were absent and neither powder was magnetic, confirming that reaction between iron and TiAl 3 had occurred. Several peaks emerged which could not be assigned to specific phases in the PDF database [43] and could represent a new phase. However, detailed examination of this is not appropriate in this work and more detailed examination must be performed before this can be confirmed.
5. Conclusions Milling of a mixture of ilmenite and aluminium lead to reaction within the mill forming TiAl 3 , Fe 4 Al 13 and Al 2 O 3 . The same reaction was found to occur in unmilled powders after heating to .8508C for 1 h, but the extent of reaction was lower. Although no distinct crystalline phases were observed by XRD directly after milling, isothermal annealing at 8008C caused recrystallisation enabling detection. Differential thermal analysis showed that the aluminium had been completely consumed below its melting point in the milled powder, whereas it was abundant in the unmilled powder. The reaction seemed to occur in two stages, formation of elemental iron and a titanium phase followed by a slow diffusion controlled reaction between iron and aluminium. In the milled powder there was no evidence for elemental iron indicating that both stages of reaction occurred, whereas the unmilled powder became magnetic on annealing for 1 h at up to 10008C. The only thermal process occurring in the milled powder was a phase transition at |10708C, which became present in the unmilled powder only after annealing for a few hours. A final slow reaction, between iron and TiAl 3 occurred during extended annealing at high temperature forming an unidentified phase.
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References [1] J.H. Westbrook, MRS Bull. 21 (1996) 37. [2] H. Morimoto, M. Nomura, Y. Ashida, J. Jpn. Inst. Met. 59 (1995) 429. [3] A.V. Smith, D.D.L. Chung, J. Mater. Sci. 31 (1996) 5961. [4] S. Linderoth, M.S. Pedersen, J. Appl. Phys. 75 (1994) 5867. [5] H.C. Ryu, M.K. Shim, S.H. Hong, J. Mater. Process. Technol. 63 (1997) 411. [6] D. Osso, A. Mocellin, G. Lecaer, A. Pianelli, J. de Physique IV 3 (1993) 1311. [7] M.D. Alcala, P. Matteazzi, J. Radioanalytical Nucl. Chem. 190 (1995) 341. [8] N. Claussen, D.E. Garcia, R. Janssen, J. Mater. Res. 11 (1996) 2884. [9] N.J. Welham, Intermetallics, (1998) in press. [10] D.P. Pope, R. Darolia, MRS Bull. 21 (1996) 30. [11] L.M. Schetky, MRS Bull. 21 (1996) 50. [12] N.J. Welham, Minerals Engrg. 9 (1996) 1189. [13] N.J. Welham, D.J. Llewellyn, (1998) submitted to Minerals Engineering. [14] A. Roine, HSC Chemistry for Windows, 2.0, Outokumpu Research Oy, Pori, 1994. [15] D.D. Wagman, W.H. Evans, V.B. Parker, I. Halow, S.M. Bailey, R.H. Schumm, Standard Values of Chemical Thermodynamic Properties, Technical Note 270-4, NBS, Washington, 1969, p. 74. [16] A. Calka, A.P. Radlinski, Mater. Sci. Engrg. A 134 (1991) 1350. [17] W.W. Wendtland, Thermal Methods of Analysis, Interscience, New York, 1964, p. 424. [18] M.I. Pope, M.D. Judd, Differential Thermal Analysis, Heyden, London, 1977, p. 197. [19] D. Dollimore, M. Reading, Application of thermal analysis to kinetic evaluation of thermal decomposition, in: J.D. Winefordner (Ed.), Treatise on Analytical Chemistry, Interscience, New York, 1993, p. 1. [20] D. Dollimore, Anal. Chem. 66 (1994) 17R. [21] B.E. Warren, X-Ray Diffraction, Dover, New York, 1990, 251. [22] D.R. Lide, Handbook of Chemistry and Physics, CRC Press, Boca Raton, 1992, 5.92. [23] J.H. Sharp, Reaction Kinetics, in: R.C. Mackenzie (Ed.), Differential Thermal Analysis, Academic Press, London, 1972, p. 47. [24] X. Wang, H.Y. Sohn, M.E. Schlesinger, Mater. Sci. Engrg. A Struct. Mater. Properties Microstructure Process. 186 (1994) 151. [25] H.Y. Sohn, X. Wang, J. Mater. Sci. 31 (1996) 3281. [26] W.B. Hillig, Discussion, Chapters 32–35, in: W.D. Kingery (Ed.), Kinetics of High-Temperature Processes, Wiley, New York, 1959, 311. [27] R.G. Becher, R.G. Canning, B.A. Goodheart, S. Uusna, Proc. of the Australasian Institute of Mining and Metallurgy, 214 (1965) 21. [28] I.E. Grey, D.G. Jones, A.F. Reid, Trans. Inst. Mining Metall. 82C (1973) 151. [29] D.G. Jones, Trans. Inst. Mining Metall. 82C (1973) 186. [30] I.E. Grey, A.F. Reid, Trans. Inst. Mining Metall. 83C (1974) 39. [31] D.G. Jones, Trans. Inst. Mining Metall. 83C (1974) 1. [32] U. Schwertmann, J. Friedi, G. Pfab, A.U. Gehring, Clays Clay Minerals 43 (1995) 599. [33] N.J. Welham, T. Hwang, M.H. Marsh, unpublished results, 1996. [34] N.J. Welham, Trans. Inst. Mining Metall., Section C. Mineral Process. Extractive Metall. (1998) in press. [35] A. Kerr, N.J. Welham, P.E. Willis, Nanostructured Mater. (1998) in press. [36] N.J. Welham, A. Kerr, P.E. Willis, J. Am. Ceramic Soc. (1998) in press. [37] P.E. Willis, N.J. Welham, A. Kerr, J. Eur. Ceramic Soc. (1998) in press. [38] R. Hultgren, R.L. Orr, P.D. Anderson, K.K. Kelley, Selected Values of Thermodynamic Properties of Metals and Alloys, Wiley, New York, 1963, 414.
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[39] T.B. Massalaki, H. Okamoto, P.R. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, Metals Park, Ohio, 1990. [40] J.C. Schuster, H. Ipser, Zeitschrift fur Metallkunde 81 (1990) 389. [41] U.R. Kattner, J.C. Lin, Y.A. Chang, Metall. Trans. A 23 (1992) 2081.
[42] D. Dew-Hughes, L. Kaufman, CALPHAD 3 (1979) 175. [43] ICDD, Powder Diffraction File, 1996, up to set 46. [44] T.I. Yanson, N.B. Manyako, O.I. Bodak, R. Cerny, R.E. Gladyshevskii, K. Yvon, J. Alloys Comp. 219 (1995) 135.
L
Mechanochemical reaction between ilmenite (FeTiO 3 ) and aluminium N.J. Welham* Petrochemistry and Experimental Petrology, Research School of Earth Sciences, and Department of Electronic Materials Engineering, Research School of Physical Sciences and Engineering, Australian National University, Canberra ACT 0200, Australia Received 2 December 1997
Abstract A natural ilmenite (FeTiO 3 ) and aluminium powder have been mechanically milled together for 100 h in a laboratory ball mill. The as-milled powder and an unmilled powder of identical composition were annealed at up to 12008C and examined by X-ray diffraction and differential thermal analysis (DTA). The unmilled sample showed aluminium melted prior to an exothermic reaction starting at |8508C. The milled powder showed no thermal activity, other than a reversible phase transition at 1067648C, indicating that reaction occurred within the mill. The products of both powders were the same, TiAl 3 , Fe 4 Al 13 and Al 2 O 3 , although in the milled powder these phases were nanocrystalline until annealing caused crystallite growth. The thermal reaction seemed to occur in two stages, formation of TiAl 3 , Al 2 O 3 and elemental iron followed by a slower, diffusion controlled reaction between the elemental iron and residual aluminium to form Fe 4 Al 13 . The reaction during milling was attributed to increased intermixing between the ilmenite and aluminium causing a change in the rate determining step from solid-state diffusion to another, unknown, controlling mechanism. 1998 Elsevier Science S.A. Keywords: Aluminium metal-oxide systems; Mechanical milling; Solid state diffusion
1. Introduction There has been increasing interest in the properties and uses of intermetallic compounds over the past few years, primarily due to their superior properties when compared with those of the constituent metals. It is well known that dispersion of a hard material e.g. SiC, Al 2 O 3 within the metal matrix can give enhanced mechanical properties over that of the pure intermetallic compound. Particles are often used as the hard phase, although fibres are more common, giving strength improvements of 30% in some cases [1], the particles can be formed by in situ thermal reaction [2] or added during a further processing stage such as infiltration [3] or milling [4,5]. Several aluminium–metal oxide systems have been milled, leading to reduction within the mill to form an alumina dispersion in the metal [6–8]. A previous paper [9] has indicated that the reaction between aluminium and TiO 2 to form alumina dispersed in TiAl 3 has its thermal onset temperature lowered by 5008C by milling the powders prior to heating. Additionally, the reaction became potentially self-propagating after milling, whereas it was net endothermic without milling. There may also be advantages in using mixtures of *Tel.: 0061 2 6249 0520; fax: 0061 2 6249 0511; e-mail: [email protected] 0925-8388 / 98 / $19.00 1998 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 98 )00123-6
intermetallic compounds [1,10,11], although the specific gains will vary with the combination used. This paper examines the milling of a mixture of naturally occurring ilmenite and aluminium to determine whether a composite of alumina dispersed in a mixture of iron and titanium aluminides can be formed at low temperature.
2. Experimental The ilmenite concentrate was the same as that detailed previously [12] and contained ,5% impurities. The asreceived particle size was 90% 210690 mm, but the powder was milled in vacuum for 25 h prior to use to reduce the particle size to .95% ,2 mm [13]. The aluminium powder was nominally .99% Al, although some surface oxide was evident, the particle size was ,20 mm. Feed powders were prepared in accordance with the stoichiometry of reaction (1). This stoichiometry was chosen on the basis of thermodynamic predictions, using free energy minimisation, which showed 7:1 to be the ratio beyond which additional aluminium remained unreacted. However, this prediction was made using species for which thermodynamic data was available in the database [14] with the addition of data at 298.15 K for the iron
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aluminides FeAl, FeAl 2 and FeAl 3 [15]. Since no data was found for the phases Fe 3 Al, Fe 2 Al 5 and Fe 4 Al 13 this reaction can only be considered a guide. FeTiO 3 1 7Al ⇒ FeAl 2 1 TiAl 3 1 2Al 2 O 3
(1)
Duplicate samples of each mixture (with a nominal 10 mol % excess of aluminium to account for surface oxide) were sealed in separate laboratory-scale ball mills. Five 10 (25.4 mm) diameter stainless steel balls were added to one mill giving a ball:powder mass ratio of 47:1. The mills were evacuated to ¯10 Pa and rotated at 165 rpm for 100 h, using magnets to control ball motion within the mill [16]. After milling, there was no change in gas pressure, indicating that gas was neither evolved nor leaked into the system. The sample derived from the mill which contained balls was considered to be milled, that from the mill with no balls was unmilled. Differential thermal analysis (DTA) was performed on the milled and unmilled powders. Samples of 3062 mg were loaded into an alumina crucible and heated up in an inert argon atmosphere to 12008C and cooled to 5008C, both at 208C min 21 using a Shimadzu DTA-50 instrument. Isothermal experiments were also performed on the unmilled powder to determine the rate controlling step by modelling the DTA data [17–20]. For the optimised model at each temperature, values of the rate constant were plotted against reciprocal temperature and activation energy of reaction determined. Thermogravimetric analysis (TGA) was performed under identical conditions in a Shimadzu TGA-50. Samples of powder were also isothermally annealed under flowing argon for 1 h. The products were analysed by X-ray diffraction (XRD) ˚ using monochromatised Co Ka radiation ( l 51.78896 A) using a count time of 2 s per 0.018 step over the range 10–1108. A narrower range of 20–708 is presented in the figures to aid clarity, but calculations and identification were made using the complete range of 2u measured. The crystallite size was estimated by applying the Scherrer equation [21] to all of the available identified peaks for each phase. The error reported in the crystallite size is 2s for the available data. Scanning electron microscopy was used to determine morphology and the distribution of the iron, aluminium and titanium within selected powders.
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shown in Fig. 1(a), no thermal reactions were detected below this temperature for any powder. During the heating of the unmilled powder there is a large endotherm around 6608C, this is due to the melting of the aluminium in the powder. The energy associated with this melting was equivalent to 0.36 kJ / g Al, which is 10% lower than the accepted value of the enthalpy of fusion for Al (0.397 kJ / g Al) [22]. This would indicate that |10% of the mass of the aluminium powder was present as oxide, indicating that the apparent 10% molar excess of aluminium used should give almost the exact stoichiometry for reaction (1). After the aluminium had melted, there was an exotherm which started at 8658C and was completed by 10908C. The energy associated with this exotherm was 0.32 J / g of powder (0.55 kJ / g Al) which is considerably less than the 1.0 kJ / g necessary [14] to heat the mixture of ilmenite and aluminium up to the ignition point at 8658C indicating that the reaction cannot be self-propagating. A slower heating rate of 58C min 21 showed that the peak was due to a single reaction. On cooling the system, the only event was the freezing
3. Results
3.1. Unmilled powder TGA of unmilled powder showed that there were no mass changes (,0.5%) when heated up to 12008C, confirming the entirely solid state nature of the reactions occurring, as expected from reaction (1). DTA traces of the unmilled powder for T .5008C is
Fig. 1. DTA traces, normalised to equal mass, for powders heated to 12008C at 208C min 21 and cooled at 208C min 21 to 5008C under argon. The dashed lines represent the cooling portion of the cycle. (a) unmilled, first cycle; (b) unmilled, second cycle and (c) milled for 100 h.
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of residual aluminium, the energy evolved indicated that |24% of the aluminium had not reacted during the heating excursion. Trace (b) in Fig. 1 shows the results of heating the residue of trace (a) up to 12008C again. Clearly, the only major thermal event was the melting and freezing of the unreacted aluminium. The energy associated with the freezing was only 85% that of the melting suggesting that aluminium was being consumed by a reaction above the melting point, this was confirmed by subsequent heating / cooling cycles which showed a slow decrease in the peaks associated with elemental aluminium. The absence of a clear thermal event involving aluminium was most probably due to it occurring slowly, giving a rise in the differential temperature which was similar to that of the background noise. This would suggest that reduction within the system was completed during the initial heating or, an exotherm associated with alumina formation would be expected due to its large enthalpy of formation from aluminium (21676 kJ mol 21 at 298 K). The slowness of this second reaction would explain the diminution of residual aluminium with the number of heating excursions up to 12008C. Indeed, DTA of a powder annealed at 12008C for 1 h showed a similar trace to that of Fig. 1(c), but with the addition of a small pair of peaks for melting / freezing of aluminium. Clearly, even 1 h at 12008C was insufficient for complete reaction to occur, although the annealed powder was becoming similar to that of the powder milled for 100 h. Isothermal DTA was performed on the unmilled powder at several temperatures within the exotherm. Modelling [17–20,23] of the exotherm showed that an Avrami– Erofeev model with exponent 0.67 was the best fit, as demonstrated by the plot inset in Fig. 2. The physical meaning of this model is that the reaction rate is controlled by the movement of the reaction interface in three dimensions [23]. The Arrhenius plot, Fig. 2, shows that the activation energy for the reaction was 358 kJ mol 21 , somewhat lower than the values of 517 and 483 kJ mol 21 found for the formation of TiAl 3 from elemental aluminium and titanium [24,25]. The latter of these values was obtained using a second order reaction model, which has been considered to have no physical meaning in solid-state reactions [23,26] and, hence is open to uncertainty. The onset temperature for the reduction step (8808C) was somewhat greater than that for synthesis from the elements (7308C) [24]. The clear implication is that the controlling step of reaction (1) was not formation from the elements, as would be expected since thermal reduction of TiO 2 by excess Al did not show any evidence for the formation of elemental titanium [9]. Examination of the XRD trace of the unmilled powder, shown in Fig. 3(a), indicates that aluminium and ilmenite are the only phases present. The aluminium shows narrow peaks indicating a large crystallite size (7366 nm), whereas the ilmenite shows somewhat broader, less intense peaks typical of a phase of small crystallite size (17610
Fig. 2. Arrhenius plot of the rate constant for the data from unmilled powder isothermally heated in the DTA. The rate constant was derived from the mathematical model of da / dt versus t based on the Avrami– Erofeev equation with exponent 0.67. A sample plot of the model, for 9508C, is shown inset, with the hollow squares the data points and the line the model, the data points shown are 20% of those used in the fitting to aid clarity.
nm). This refinement of ilmenite crystallite size was due to the 25 h premilling prior to mixing. A sample of the unmilled powder was annealed at 6008C for 1 h, the XRD trace of this powder is shown in Fig. 3(b). There were no new peaks present, indicating that no apparent reaction had occurred. The only change was a narrowing of the peaks due to thermally induced crystallite growth, giving 2868 nm and 8869 nm for ilmenite and aluminium respectively. A broad peak around 318 was also evident, this was probably due to rutile (TiO 2 ). The initial ilmenite powder contained a small amount of rutile [12] and its presence may be due to crystallite growth, which has been noted to occur below 6008C [9]. Rutile is also the initial reduction product of ilmenite [27–31] and its presence could also be due to reduction. The other product of ilmenite reduction is elemental iron, unfortunately, the main peaks for elemental iron are coangular with peaks of aluminium and cannot be deconvoluted, primarily due to the low intensity of the iron peaks that would be expected
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Fig. 3. XRD traces for powders: (a) unmilled; (b) unmilled, 1 h at 6008C; (c) unmilled, 1 h at 10008C; (d) 100 h milled; (e) 100 h milled, 1 h at 6008C; (f) 100 h milled, 1 h at 10008C; (g) 100 h milled, 12 h at 12008C; m-Al, ♦-Al 2 O 3 , j-TiAl 3 , d-Fe 4 Al 13 , n-FeTiO 3 , s-Fe and h-TiO 2 data for Fe 4 Al 13 is only available for 2u #55.038.
to be present. However, the powder annealed at 6008C proved to be magnetic, the unmilled powder having showed no signs of magnetism. Clearly, some reduction reaction had occurred, presumably forming elemental iron, although the extent was probably small as there were no traces of an exotherm which would be expected when Al 2 O 3 forms from Al. The reaction occurred below the melting point of aluminium and can only have occurred at the contact points between the large aluminium grains and the smaller ilmenite particles. Annealing above the melting point of aluminium but below the exotherm, also gave a magnetic powder, XRD of which showed weak peaks where the main peaks of alumina and TiAl 3 were expected in addition to those of aluminium and ilmenite. Clearly, the melting of aluminium causes an increase in contact area between the reactant phases leading to a larger extent of reaction which allows detection by XRD. Only after annealing above the exotherm, were major phase changes observed in the XRD trace, as in trace (c). Both TiAl 3 and Al 2 O 3 emerged, the main aluminium peak
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had weakened considerably and ilmenite was absent, thus confirming that reaction (1) had occurred to some extent. The powder was magnetic, implying that elemental iron was still present, but its XRD peaks concealed. The rutile present at 6008C had been consumed by reduction. There were several peaks which were probably due to Fe 4 Al 13 , the most intense of these appearing as shoulders on the aluminium and alumina peaks at 51 and 538. The continued presence of aluminium after 1 h of annealing above the exotherm is a clear indicator that the formation of one phase is slow. The evidence seems to be that TiAl 3 is readily formed, with clear peaks present after 1 h at 10008C, Fig. 3(c). However, no iron bearing phase can be confirmed as present, although peaks are probably present for Fe 4 Al 13 . This would imply that the final iron bearing phase was the one which formed slowly. The free energy of formation of TiAl 3 from its elements is 2170 kJ mol 21 whereas that for FeAl 3 is 279 kJ mol 21 (the closest phase to Fe 4 Al 13 for which any data was available), clearly the formation of TiAl 3 is more favourable. The reduction sequence for ilmenite passes through elemental iron and TiO 2 [27–31] and it may be expected that the aluminium would react with the elemental iron in preference to the TiO 2 . However, the free energy for the reduction of TiO 2 by Al to alumina and titanium trialuminide is 2301 kJ mol 21 TiO 2 —much more favourable than reaction of aluminium with iron to form iron aluminide. There was no XRD evidence for intermediate phases, even after quenching DTA runs at points on the reaction exotherm corresponding to |10, 25 and 50% aluminium consumption. The magnetism evident in the annealed powders would seem to indicate that elemental iron was formed during the reduction and remained in the powder after 1 h at 10008C. Thus, it would seem that the rate determining step is the reaction of aluminium with the titanium phase from ilmenite reduction. Previous work [9] has shown that the thermally induced reaction between a mixture of aluminium and rutile, prepared identically to that used here, occurs at 10008C. Thus, it appears that the titanium product of ilmenite reduction was not crystalline rutile, or the main reaction exotherm would have been observed at higher temperature. The titanium phase formed may have been nanocrystalline and thus, not readily detectable by XRD. The separation of iron and titanium at the atomic level in ilmenite could cause more intimate mixing between the titanium phase and aluminium than in physical mixtures, lowering the reaction temperature as supported by the onset temperature of 6008C for milled rutile and aluminium [9]. The presence of elemental iron may also act as a catalyst by forming intermediate phases which reduces the energetic requirements. It has been shown that Fe 2 TiO 5 forms by reaction between elemental iron and TiO 2 at temperatures above 6008C [32,33] and phases of this type may be more amenable to reduction than TiO 2 . Regardless of the nature of the initial rate controlling
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step, the temperature required for this step to occur is greater than that necessary for the subsequent reactions forming alumina and TiAl 3 which consequently occur extremely rapidly. The final iron phase, Fe 4 Al 13 , would seem to be the phase formed relatively slowly as it was only detected after annealing for 1 h at 10008C. Its formation is presumably by interdiffusion of elemental iron and aluminium which would be expected to be slow at low temperatures giving a phase of varying composition and low crystallinity, which would be difficult to detect by XRD. With increased annealing temperature, the quantity of Fe 4 Al 13 formed increased and crystallite growth occurred, rendering it more readily detectable by XRD. Indeed, after 12 h at 12008C the powder was no longer magnetic showing that elemental iron had been completely consumed by reaction, no peaks for aluminium were evident in the DTA. The slow diminution of the aluminium melting / freezing peaks in the DTA with thermal cycling would support the slow consumption of aluminium, the overlap of the XRD peaks for iron and aluminium concealing the concomitant consumption of elemental iron.
3.2. Milled powder The DTA trace for the milled powder is shown in Fig. 1(c), the complete absence of any aluminium was clear, with no evidence for a melting endotherm. This proves that reactions involving elemental aluminium were completed within the mill. The only thermal events observed were the endotherm and exotherm of a phase change around 10708C. It is interesting to note the difference in shape between the heating and cooling events. The heating shows a symmetrical peak, but on cooling there are two overlapping peaks. The energy given out during cooling was around 95% that taken in during heating, this ratio was constant within error for all samples. A slower heating rate (58C min 21 ) was used to determine whether the peaks were separate, this showed essentially the same result—a symmetrical peak during heating and an asymmetric cooling peak. A sample of milled powder annealed at 12008C for 1 h showed an identical DTA trace to that presented in (c). The DTA trace of a sample of powder heated up to 10508C, just below the onset of the peak during heating did not show any peaks on cooling, whereas a sample heated to 10708C, just above the onset showed an asymmetrical cooling peak. This result confirms that the asymmetry was associated with the phase change, however, the reason for this is uncertain. It is possible that there are two phases transitions at slightly different temperatures, however the absence of separate peaks during heating would seem to discount this. Separation of the two overlapping exotherms observed during cooling was made assuming both were symmetrical, the energy associated with the second peak was 30% that of the main freezing peak. The transition point of the phase was estimated as the
mean of the temperatures at the maxima of the melting and freezing peaks [17,18,20], the average of four separate experimental runs gave a value of 1067648C. The XRD trace for the milled powder, Fig. 3(d), shows a broad peak at 468 and a very broad peak between 48–558, the first of these peaks was assigned to TiAl 3 , the absence of any other significant peaks makes this tentative. The second peak was too broad for profile fitting to reveal any profitable information, although the main peaks of several possible phases (e.g. Al 2 O 3 , FeAl, FeAl 2 , FeAl 3 , Fe 4 Al 13 ) are around this area and a combination of some, or all, of these may be present. The estimated crystallite size for the peak at 468 was |6 nm, but this is subject to potentially gross error due to uncertainties in whether this peak was for a single phase or was due to multiple phases with overlapping peaks. There was no evidence for either aluminium or ilmenite, although this may simply be due to reduction of the crystallite size, or amorphisation during milling. After annealing the milled powder at 6008C the trace (e) was obtained, crystallite growth is clearly evident for the main peak and is almost certainly due to TiAl 3 . Smaller peaks are now present for alumina but other peaks cannot be attributed with any certainty, particularly around 50– 558 where a broad peak is evident. The absence of peaks for ilmenite and aluminium can now be definitely attributed to reaction since the unmilled powder (b) showed crystallite growth occurred for both of these phases at 6008C. The XRD trace for the 1 h of annealing at 10008C is shown in trace (f), in which the grain growth of several phases is evident. TiAl 3 and Al 2 O 3 both show several distinct peaks thus confirming their presence. The majority of the remaining peaks were due to Fe 4 Al 13 , the main peaks for this phase are around 51–538 and these were almost certainly contributing to the broad peak noted in (d) and (e). The confirmation of the presence of Fe 4 Al 13 in the milled sample annealed to 10008C, would indicate that it was probably also present in the unmilled sample (c). The somewhat weaker peaks for the unmilled powder are due to the low fraction of the phase present, due to solid-state diffusion limitation of its formation. Unlike the unmilled powders no magnetic material was present in the as-milled powder, nor in any of the annealed powders. Clearly, milling performed the same function as high temperature annealing of unmilled powders, giving complete reaction between ilmenite and aluminium directly after milling.
3.3. Electron microscopy The secondary electron images and elemental maps for unmilled and milled powders are shown in Fig. 4(a) and (b) respectively. The unmilled powder consisted of large rounded grains surrounded by smaller submicron grains. The elemental maps confirm that the large grains were aluminium and the small grains ilmenite, as expected from
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Fig. 4. Micrographs (scale bar is 20 mm) and their corresponding digitally collected elemental maps of titanium, iron and aluminium for powders: (a) unmilled; (b) 100 h milled; (c) unmilled, 1 h at 10008C; (d) 100 h milled, 1 h at 10008C.
the feed powder particle size. The two phases were essentially discrete, although the ilmenite did appear to form a layer on the surface of the aluminium grains, due to adhesion between the particles during tumbling in the mill. The milled powder showed somewhat different morphology, the smallest grains are not evident, giving a minimum particle size around 1 mm, although the majority are 5–15 mm. The grains were comparatively angular consisting of submicron fragments which appear to have been welded together, forming larger agglomerates. An equilibrium particle size has been shown to be a typical feature of extended milling [13,34] although different materials / mixtures attain different particle sizes. It has been noted that the presence of a metallic phase, either added initially or formed by reaction during milling, tends to give a larger particle size than in a metal free system [35–37] indicating that the metal acts as a binder for the submicron particles. Elemental maps show that the aluminium, titanium and
iron were present in all grains, showing that milling produces intimate intermixing beyond the resolution of the mapping (512 mapping points per 61 mm giving an apparent resolution of 0.12 mm, but probably closer to |0.5 mm due to the interaction volume of the beam). The reaction that took place during milling could be expected to lead to this type of intermixing, although intermixing on a submicron scale has been demonstrated previously in a system where reaction did not occur during milling [9]. After annealing, the unmilled powder showed no significant change in particle size, Fig. 4(c), with large grains and small grains still evident. The large grains appear to be more angular and faceted than those of the unheated powder (a), this is probably due to crystallisation of aluminium from the liquid phase during cooling from 10008C. Elemental mapping shows that aluminium was also present in some of the small grains, indicating that reaction had caused some breakage of the large grains. Titanium and iron were still dispersed throughout the
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powder, although several grains show separate areas for titanium and iron, indicating some segregation of iron and titanium had occurred. There are a few areas where aluminium seems to be associated with the titanium as opposed to evenly distributed between the iron and titanium. This may explain why the formation of Fe 4 Al 13 was slow, as segregated areas of aluminium and iron would increase the diffusional pathlength, compared with the well mixed areas. The micrographs and elemental maps for the milled powder annealed at 10008C for 1 h are shown in Fig. 4(d). There seems to have been an increase in the maximum particle size compared with the unannealed powders (b) and the particles appear to be more angular. The elemental maps showed the expected result, with no evidence of segregation of iron, titanium and aluminium.
4. Discussion It is clear that the same reaction occurred during mechanical milling as during thermal treatment of unmilled power of identical composition. This could imply that the instantaneous temperature during impact reached that necessary for reaction i.e. .8508C. However, modelling of the thermally induced reaction for the unmilled powder indicated that the activation energy was 358 kJ mol 21 , typical of solid-state diffusional control. Elemental mapping of the powders showed that milling caused intimate mixing of the phases, not evident in the unmilled powder. Clearly, milling must have decreased the diffusional pathlength, hence causing a change in the rate controlling step. However, the absence of thermal reaction in the milled powder does not allow the activation energy of reaction for the milled powder to be determined. In both the milled and unmilled powders the final products after annealing were the same, TiAl 3 , Fe 4 Al 13 and alumina, although the extent of reaction was greater for the milled powder than the unmilled with residual aluminium in the unmilled powder. The thermal reaction seemed to occur in two main stages, the formation of alumina, TiAl 3 and elemental iron, followed by a slow diffusion controlled reaction between residual aluminium and iron to form Fe 4 Al 13 . A third reaction between elemental iron and TiAl 3 was also evident, but only after extensive annealing at high temperature. For the milled powder, the main reactions occurred within the mill, there being no residual aluminium detected in the DTA during initial heating. For milled samples pretreated in different ways, the only thermal process detected was a reversible phase transition at around 10678C, therefore elemental aluminium was definitely absent. Examination of the phase diagrams of the binary systems for the non-oxides that were present [38–41] showed no phases had melting points close to this value. A ternary diagram for the Fe–Al–Ti system [42] did not show any evidence for a three
component phase with a phase transition, although at least two further phases within this system (Ti 0.75 Fe 0.25 Al 3 and Fe 4 Ti 0.93 Al 12.07 ) have since been crystallographically described [43,44], but were not present. Examination by XRD provided no clues as to the identity of this phase, with only phases of higher melting point detected. A sample of unmilled powder, annealed for 1 h at 12008C showed similar DTA behaviour to that of the as-milled powder, with a transition point of |10708C evident. The presence of Fe 4 Al 11 would seem to be necessary for this energetic event to be detected. Phase diagrams for the iron–aluminium system [38,39] indicate a melting point around 11508C for a phase of this composition, but no phase transition has been reported. The overlapping peaks observed on cooling in the DTA may be due to separate phases of similar composition. During heating, the phases undergo transformation at a similar temperature that cannot be separated by DTA. Above the transition temperature, a diffusion reaction occurs which changes the relative stoichiometry of the phases slightly, but insufficient to change the shape of the DTA trace on cooling, one phase reversed its transition at a slightly higher temperature than the other giving rise to two peaks. A sample of milled powder was annealed at 12008C for 12 h, DTA on this sample showed a transition point at |10708C that did not show asymmetry on cooling, clearly, this long annealing time was sufficient for homogenisation to occur, giving a single uniform stoichiometry. The original stoichiometry used contains insufficient aluminium for complete formation of Fe 4 Al 13 , TiAl 3 and Al 2 O 3 . A ratio of 8.25 Al per ilmenite is required for complete reaction, whereas the amount used was |7 Al per ilmenite (accounting for 10% of oxide). Assuming that TiAl 3 and alumina form and all remaining Al forms Fe 4 Al 13 then 38.5% of the iron should be present as metal. This should be readily distinguishable from aluminium by the absence of a peak at |458 whilst the peaks at larger 2u should remain, however, XRD data for Fe 4 Al 13 is only available for 2u #55.038 and overlap may occur. A sample of powder which contained Fe 4 Al 13 did not show peaks at the larger 2u peaks of iron, thus confirming that the higher 2u peaks did not overlap and were diagnostic for iron. Iron was certainly present in the unmilled powder products as they were all magnetic, none of the milled powders were magnetic showing that, although there was excess iron in the system, it has been rendered nonmagnetic, presumably by reaction with aluminium to form an aluminide of uncertain composition. This iron aluminide could be the reason for the second peak during cooling of the DTA. The available thermodynamic data indicates that iron may react with the higher aluminides to form less aluminium rich aluminides, e.g. by reaction (2) or (3) which have respective free energies of 271 and 213 kJ mol 21 Fe at 258C. Similar calculations for reaction between titanium tri-aluminide and iron showed that reaction (4) was only just favourable at 258C with a free energy of 21
N. J. Welham / Journal of Alloys and Compounds 270 (1998) 228 – 236
kJ mol Fe, and hence would probably be a minor reaction compared with (2) and (3). The free energy at higher temperatures, where interdiffusion would be significant, could not be calculated due to the absence of entropy and heat capacity data for the iron aluminides in the data source used [15]. 2FeAl 3 1 Fe ⇒ 3FeAl 2
(2)
FeAl 2 1 Fe ⇒ 2FeAl
(3)
2Fe 1 TiAl 3 ⇒ TiAl 1 2FeAl
(4)
Milled and unmilled powders were annealed for 12 h at 12008C under argon to assess the possibility of reactions of these type occurring. The XRD traces of the resultant powders were essentially identical, with only a slight change in the relative intensities of the peaks. The milled powder trace, Fig. 3(g), showed only a weak TiAl 3 peak, the peaks for Al 2 O 3 and Fe 4 Al 13 remained approximately constant. Peaks at higher 2u which correspond to iron (and aluminium) were absent and neither powder was magnetic, confirming that reaction between iron and TiAl 3 had occurred. Several peaks emerged which could not be assigned to specific phases in the PDF database [43] and could represent a new phase. However, detailed examination of this is not appropriate in this work and more detailed examination must be performed before this can be confirmed.
5. Conclusions Milling of a mixture of ilmenite and aluminium lead to reaction within the mill forming TiAl 3 , Fe 4 Al 13 and Al 2 O 3 . The same reaction was found to occur in unmilled powders after heating to .8508C for 1 h, but the extent of reaction was lower. Although no distinct crystalline phases were observed by XRD directly after milling, isothermal annealing at 8008C caused recrystallisation enabling detection. Differential thermal analysis showed that the aluminium had been completely consumed below its melting point in the milled powder, whereas it was abundant in the unmilled powder. The reaction seemed to occur in two stages, formation of elemental iron and a titanium phase followed by a slow diffusion controlled reaction between iron and aluminium. In the milled powder there was no evidence for elemental iron indicating that both stages of reaction occurred, whereas the unmilled powder became magnetic on annealing for 1 h at up to 10008C. The only thermal process occurring in the milled powder was a phase transition at |10708C, which became present in the unmilled powder only after annealing for a few hours. A final slow reaction, between iron and TiAl 3 occurred during extended annealing at high temperature forming an unidentified phase.
235
References [1] J.H. Westbrook, MRS Bull. 21 (1996) 37. [2] H. Morimoto, M. Nomura, Y. Ashida, J. Jpn. Inst. Met. 59 (1995) 429. [3] A.V. Smith, D.D.L. Chung, J. Mater. Sci. 31 (1996) 5961. [4] S. Linderoth, M.S. Pedersen, J. Appl. Phys. 75 (1994) 5867. [5] H.C. Ryu, M.K. Shim, S.H. Hong, J. Mater. Process. Technol. 63 (1997) 411. [6] D. Osso, A. Mocellin, G. Lecaer, A. Pianelli, J. de Physique IV 3 (1993) 1311. [7] M.D. Alcala, P. Matteazzi, J. Radioanalytical Nucl. Chem. 190 (1995) 341. [8] N. Claussen, D.E. Garcia, R. Janssen, J. Mater. Res. 11 (1996) 2884. [9] N.J. Welham, Intermetallics, (1998) in press. [10] D.P. Pope, R. Darolia, MRS Bull. 21 (1996) 30. [11] L.M. Schetky, MRS Bull. 21 (1996) 50. [12] N.J. Welham, Minerals Engrg. 9 (1996) 1189. [13] N.J. Welham, D.J. Llewellyn, (1998) submitted to Minerals Engineering. [14] A. Roine, HSC Chemistry for Windows, 2.0, Outokumpu Research Oy, Pori, 1994. [15] D.D. Wagman, W.H. Evans, V.B. Parker, I. Halow, S.M. Bailey, R.H. Schumm, Standard Values of Chemical Thermodynamic Properties, Technical Note 270-4, NBS, Washington, 1969, p. 74. [16] A. Calka, A.P. Radlinski, Mater. Sci. Engrg. A 134 (1991) 1350. [17] W.W. Wendtland, Thermal Methods of Analysis, Interscience, New York, 1964, p. 424. [18] M.I. Pope, M.D. Judd, Differential Thermal Analysis, Heyden, London, 1977, p. 197. [19] D. Dollimore, M. Reading, Application of thermal analysis to kinetic evaluation of thermal decomposition, in: J.D. Winefordner (Ed.), Treatise on Analytical Chemistry, Interscience, New York, 1993, p. 1. [20] D. Dollimore, Anal. Chem. 66 (1994) 17R. [21] B.E. Warren, X-Ray Diffraction, Dover, New York, 1990, 251. [22] D.R. Lide, Handbook of Chemistry and Physics, CRC Press, Boca Raton, 1992, 5.92. [23] J.H. Sharp, Reaction Kinetics, in: R.C. Mackenzie (Ed.), Differential Thermal Analysis, Academic Press, London, 1972, p. 47. [24] X. Wang, H.Y. Sohn, M.E. Schlesinger, Mater. Sci. Engrg. A Struct. Mater. Properties Microstructure Process. 186 (1994) 151. [25] H.Y. Sohn, X. Wang, J. Mater. Sci. 31 (1996) 3281. [26] W.B. Hillig, Discussion, Chapters 32–35, in: W.D. Kingery (Ed.), Kinetics of High-Temperature Processes, Wiley, New York, 1959, 311. [27] R.G. Becher, R.G. Canning, B.A. Goodheart, S. Uusna, Proc. of the Australasian Institute of Mining and Metallurgy, 214 (1965) 21. [28] I.E. Grey, D.G. Jones, A.F. Reid, Trans. Inst. Mining Metall. 82C (1973) 151. [29] D.G. Jones, Trans. Inst. Mining Metall. 82C (1973) 186. [30] I.E. Grey, A.F. Reid, Trans. Inst. Mining Metall. 83C (1974) 39. [31] D.G. Jones, Trans. Inst. Mining Metall. 83C (1974) 1. [32] U. Schwertmann, J. Friedi, G. Pfab, A.U. Gehring, Clays Clay Minerals 43 (1995) 599. [33] N.J. Welham, T. Hwang, M.H. Marsh, unpublished results, 1996. [34] N.J. Welham, Trans. Inst. Mining Metall., Section C. Mineral Process. Extractive Metall. (1998) in press. [35] A. Kerr, N.J. Welham, P.E. Willis, Nanostructured Mater. (1998) in press. [36] N.J. Welham, A. Kerr, P.E. Willis, J. Am. Ceramic Soc. (1998) in press. [37] P.E. Willis, N.J. Welham, A. Kerr, J. Eur. Ceramic Soc. (1998) in press. [38] R. Hultgren, R.L. Orr, P.D. Anderson, K.K. Kelley, Selected Values of Thermodynamic Properties of Metals and Alloys, Wiley, New York, 1963, 414.
236
N. J. Welham / Journal of Alloys and Compounds 270 (1998) 228 – 236
[39] T.B. Massalaki, H. Okamoto, P.R. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, Metals Park, Ohio, 1990. [40] J.C. Schuster, H. Ipser, Zeitschrift fur Metallkunde 81 (1990) 389. [41] U.R. Kattner, J.C. Lin, Y.A. Chang, Metall. Trans. A 23 (1992) 2081.
[42] D. Dew-Hughes, L. Kaufman, CALPHAD 3 (1979) 175. [43] ICDD, Powder Diffraction File, 1996, up to set 46. [44] T.I. Yanson, N.B. Manyako, O.I. Bodak, R. Cerny, R.E. Gladyshevskii, K. Yvon, J. Alloys Comp. 219 (1995) 135.