In an attempt to reduce consumption of raw materials and wastage, recycling is becoming an ever important issue.

Although the steel industry recycles a good deal of scrap material, the amount usable is limited by the accumulation of undesirable elements within the steel. With the introduction of new EU targets for the recycling of 250,000 tonnes of tin plate scrap per year in the UK alone, it is envisaged that levels of undesirable elements will soon rise to unacceptable levels. For example, as a result of the increase in the amount of recycled tin plate, the residual tin content in basic oxygen steel making will rise from currently acceptable levels of around 0.006wt% to unacceptable levels of 0.05wt% or more.

The harmful effect of low levels of copper and/or tin on mechanical properties of steel has been recognised for some time. The levels of tin and copper resulting from the reuse of tin plate and incinerator scrap have to be carefully controlled because of the effect they have on the hot formability of steel. The segregation of tin in steel is well documented elsewhere. The presence of elemental copper in (for example) incinerator scrap leads to the precipitation of copper as a film at the grain boundaries.

Hot working of the material causes the copper to melt resulting in surface cracking and loss of ductility in the temperature range 1050 to 1200°C; an effect known as"hot shortness"which leads to surface defects on the worked material. The presence of tin is known to exacerbate the effect of the copper by forming a tin-copper alloy with a lower melting point. Generally speaking, the aim should be to keep the level of copper below around 0.2wt% and tin below about 0.05wt% to prevent precipitation at grain boundaries. Traditionally, tin plate scrap is mixed with other higher quality grade scrap in order to dilute the amount of tin.

De-tinning plants aim to reduce the level of tin by electrochemical or chemical methods. These methods are generally costly in financial and environmental terms.

Moreover, existing de-tinning plants-have limited capacity and environmental considerations place a constraint on expansion.

The present invention is based on the recognition that if an undesirable element (eg tin or copper) is rendered innocuous, increasing levels of the element in recycled steel may be controlled. The invention seeks to provide a simple, cheap method for rendering the tin and copper innocuous within the iron-rich matrix by the addition of an appropriate alloying metal to provide steel with improved hot working and mechanical properties. It will be appreciated that a method based on innocuating existing levels of an element by adding a metal is a radical departure from traditional methods which aim to reduce the level of the element by extraction.

Thus viewed from one aspect the present invention provides a process for innocuating at least one element selected from the group consisting of copper and tin present in steel, said process comprising: adding to the steel a source of an alloying metal, wherein said alloying metal exhibits a large negative enthalpy of mixing with iron in the liquid phase and a higher (eg a less negative or a positive) enthalpy of mixing in the liquid phase with the element which it is intended to innocuate (ie copper or tin).

Preferably the minimum of the enthalpy of mixing with iron in the liquid phase is in the range 0 to-60kJ/mol, preferably-10 to-40kJ/mol, particularly preferably-15 to -25kJ/mol, especially preferably about-20kJ/mol.

Preferably the enthalpy of mixing of the alloying element with copper or tin in the liquid phase is greater than the minimum enthalpy of mixing with iron in the liquid phase.

In a preferred embodiment, the alloying metal forms or causes the formation of copper and/or tin-containing phases thereby displacing copper or tin from the iron-rich matrix or grain boundary sites.

In a preferred embodiment, the alloying metal is one which induces a change in the morphology of copper in the iron- rich matrix or at grain boundaries in the steel.

In accordance with the invention, the alloying metal is one which exhibits a highly negative enthalpy of mixing with iron in the liquid phase. This is indicative of a strong attraction between the alloying metal and iron.

Additionally, the alloying metal may exhibit a small negative or a positive enthalpy of mixing in the liquid phase with the element which it is intended to innocuate (ie copper or tin) indicative of a weak attraction to or a repulsion from the copper or tin.

Preferred alloying metals are generally selected from the group consisting of Al, Ti, Si, Zr and Nb. The steel upon which the process of the invention may be carried out comprise elements other than iron and carbon in varying amounts eg Nb, Mn, Si.

The preferred alloying metal for use in the process of the invention is aluminium. The source of aluminium may be pure metal, an aluminium alloy or an aluminium compound capable of dissociation at the operating temperature. The source of aluminium typically provides a level of between 0.5 and 6wt%, preferably 1 to 5wt%, particularly preferably 3 to 4.5 wt%, especially preferably about 4wt% of aluminium metal.

In a particularly preferred embodiment, the steel upon which the process of the invention is carried out comprises a level of niobium sufficient to form the Fe2Nb-Laves phase.

Tin is known to dissolve in Fe2Nb-Laves and it has been found that addition of aluminium as the alloying metal accelerates the formation of the Fe2Nb-Laves phase and thereby increases the level of dissolved tin thereby rendering it innocuous. Preferably Nb is present in the steel upon which the process of the invention is carried out (either inherently or by addition) in the amount of about 0.9wt% or more for a steel typically containing about 0.05wt% carbon.

Not wishing to be bound by any theory, it is thought that the presence of high levels of aluminium in steel causes a change in the morphology of the copper at the grain boundary. Instead of appearing as a continuous film (see Figure 37a), copper precipitates as discrete particles along the grain boundaries (see Figure 37b).

Where high levels of copper are present, the copper tends to accumulate in the iron-rich matrix as a spherical precipitate (see Figure 36). Again not wishing to be bound by theory, it is thought that the presence of aluminium leads to a change in the precipitate morphology of copper in the matrix. Together with the presence of aluminium in the iron-rich matrix, this causes an increase-in-the hardness implying an increase in the mechanical strength (see Figure 38). More particularly, lenticular precipitation of copper shows a degree of coherency between its structure and that of the matrix.

At the present time, aluminium cans are recycled separately from steel cans (ie tin plated steel). The present invention may be carried out using mixed loads of aluminium cans and steel cans with a saving in the cost of sorting prior to use. Thus, in a preferred embodiment, the source of alloying metal used in the process of the invention is aluminium cans.

The process of the invention may be particularly useful for incinerator scrap containing up to 0.25wt% tin and 2wt% copper. It may be possible to carry out the process of the invention on mixed loads of incinerator and tin plate scrap (tin-containing steel) without the risk of producing steel of an inferior quality.

In the process of the invention, the source of alloying metal and steel are added in the liquid phase. For example, an effective amount of a source of alloying metal may be added to molten steel. This conveniently takes place in a vacuum or in an inert atmosphere. Preferably, the mixture is then heat treated at high temperature (eg about 1150°C) for 1 to 8 hours and at a lower temperature (eg about 950°C) for up to 16 hours and finally at a much lower temperature (eg about 600°C) for at least 5 hours but no more than 48 hours (eg about 16 hours).

In a preferred embodiment, the process of the invention comprises the step of obtaining the thermodynamic properties of a binary system of a candidate alloying metal and an element of the steel whereby to select candidates for innocuating the tin and/or copper. Particularly preferably, the thermodynamic property is an enthalpy of, mixing curve for the liquid phase binary system.

It has been surprisingly found that the hitherto unknown addition of high levels of aluminium to steel provides a product with a surprisingly high degree of hardness.

Viewed from a further aspect the present invention provides an aluminium containing steel with a hardness which is at least 1.3 times greater than the aluminium free analogue.

Preferably the aluminium containing steel of the invention is obtainable by a process as hereinbefore described.

Viewed from a yet further aspect the present invention provides a method for recycling steel (eg steel cans or incinerator scrap or a mixture thereof) comprising the step of innocuating copper or tin as hereinbefore described.

The present invention will now be described in a non- limitative sense with reference to the following examples and accompanying Figures in which: Fig 1 SEM micrographs of the steel S1: (a) 850°C for 96 hours and (b) 1150°C for 1 hour Fig 2 SEM micrographs of the steel S2: (a) 850°C for 96 hours and (b) 1150°C for 1 hour Fig 3 EPMA line scan analyses across precipitates in steel S1 after aging at 850°C for 96 hours: (a) Line scan of C, (b) Line scan of Nb, (c) Line scan of Si and (d) Line scan of Sn Fig 4 EPMA line scan analyses across precipitates in steel S2 after aging at 850°C for 96 hours: (a) Line scan of C, (b) Line scan of Nb, (c) Line scan of Si and (d) Line scan of Sn Fig 5 EPMA line scan analyses across precipitates in steel S2_after aging at 850°C for 96 hours: (a) Line scan of C, (b) Line scan of Nb, (c) Line scan of N and (d) Line scan of Al Fig 6 NbC formed in steel S1 after the aging at 1150°C: (a) bright field image and (b) SAD pattern of NbC from zone axis of [001] Fig 7 n-Fe2Nb3 phase formed in the lower Al steel S1 after aging at 850°C for 96 hours: (a) bright field image and (b) SAD pattern from zone axis of [lll] n Fig 8 Fe2Nb-Laves phase formed in the higher Al steel after aging at 850°C for 96 hours (a) bright field image and (b) SAD pattern of Fe2Nb-Laves phase from zone axis of [2110] Fig 9 AIN formed in the higher Al steel after aging at 850°C for 96 hours: (a) bright field image and (b) SAD pattern of AIN from zone axis of [OlT1] Fig 10 Enthalpy of mixing curves for binary liquids Fe-Al, Nb-Sn and Sl-Sn Fig 11 The compositions-of >-Fe2Nb3 and Fe2Nb-Laves phases: (a) Fe and Nb content and (b) Si and Sn content Fig 12a-d SEM micrographs of the alloys aged at 850°C for 96 hours: (a) 0.1 Nb, (b) 0.9Nb, (c) O. lTi and (d) 0.9Ti (wt%) Fig 13a-d Microprobe line scan analyses across complex- precipitates in Fe-0.5 Al-0.9Nb-0.05 Sn alloy aged at 850°C for 96 hours: (a) line scan of oxygen, (b) line scan of Al, (c) line scan of Nb and (d) line scan of tin Fig 14a-c Microprobe line scan analyses across single bright precipitates in Fe-0.5Al-O. lNb-0.05Sn alloy aged at 850°C for 96 hours: (a) line scan of Nb, (b) line scan of oxygen and (c) line scan of tin Fig 15a, b Microprobe line scan analyses across single bright precipitates in 05Sn alloy aged at 850°C for 96 hours showing the precipitate is enriched in niobium and tin: (a) line scan of Nb and (b) line scan of tin Fig 16a-d Element mapping of Al, Nb and Sn in the Fe- 0.5Al-0.9Nb-0.3Sn alloy: (a) Back scattered electron images showing complex-precipitates and single bright precipitates, (b) Element mapping of Al, (c) Element mapping of Nb and (d) Element mapping of Sn Fig 17a-d Microprobe line scan analyses precipitates in 3Sn alloy aged at 850°C for 96 hours: (a) line of scan of Al, (b) line scan of oxygen, (c) line scan of Ti and (d) line scan of tin Fig 18a, b Microprobe line scan analyses precipitates in 3Sn alloy aged at 850°C for 96 hours showing the titanium nitrides: (a) line scan of Ti and (b) line scan of nitrogen Fig 19a-d Element mapping of Al, Ti and Sn in the Fe- 3Sn alloy: (a) Back scattered electron images, (b) Element mapping of Al, (c) Element mapping of Ti and (d) Element mapping of Sn Fig 20a-d TEM studies of complex precipitates in 0.9Nb alloy: (a) Bright field micrograph, (b)- (c) SADP from zone axes of [101], [-6-11], and (d) a typical EDS spectrum of Fe2Nb3 phase Fig 21a-f TEM studies of Sn-rich precipitates in 0.9Nb alloy: (a) Bright field micrograph, (b)- (e) SADP from zone axes of [100], [1-10], [3-32] and [0-31] and (f) a typical EDS spectrum of Laves phase Fig 22a-d TEM of complex precipitates in Ti containing alloy: (a) Bright field micrograph, showing the complex precipitates (b)- (d) EDS spectra from matrix, outer part and central part of precipitates Fig 23ab TEM studies of TiN precipitates in Ti containing alloy: (a) Bright field micrograph, (b) SADP of the precipitates Figs 24-27 Results of multi-phase equilibrium calculations Fig 28 Enthalpy of mixing curves for binary liquids Fe-Al, Fe-Ti, Fe-Nb, Fe-Cu and Fe-Sn Fig 29 The Fe-Al Phase Diagram Fig 30 The Fe-Nb Phase Diagram Fig 31 Gibbs energies of mixing of the bcc phases in the Fe-Nb and Fe-Ti systems together with the Gibbs energies of the Laves phases. Tangent intercepts to the pairs of Gibbs energy curves give the respective phase boundaries Fig 32 The Fe-Sn Phase Diagram Fig 33 The Fe-Cu Phase Diagram Fig 34 Enthalpy of mixing curves for the binary liquids Nb-Sn and Al-Sn Fig 35 Enthalpy of mixing curves for the binary liquids Cu-Fe, Cu-Ti, Cu-Al and Cu-Si Fig 36 SEM of copper rich steel showing copper in the ferrite matrix (no aluminium) Fig 37 SEM of copper containing steel showing copper at the grain boundaries (a) in the absence of aluminium and (b) in the presence of aluminium Fig 38 SEM of copper rich steel containing aluminium showing lenticular precipitation of copper in the ferrite matrix Fig 39 Vickers Hardness measurements for five different heat-treated steels with and without aluminium.

Example 1 The microstructures of low carbon steels containing Sn were investigated using scanning electron microscopy, electron probe microanalysis, transmission electron microscopy and energy dispersive X-ray spectroscopy. Four steels based on Fe-0.9Nb-0.3Sn-0.05C (wt%) with different levels of Al and Si additions were prepared by arc melting under an argon atmosphere. The effects of heat treatment and level alloying elements of Al and Si on the precipitation of Sn-rich phases were studied.- Four alloys (S1 to S4) with different levels of Al and Si were prepared by arc melting under an argon atmosphere.

Repeated melting of the alloys was performed to ensure macroscopic homogeneity of the cast material. The target chemical compositions (in wt%) of the steels used in this study are listed in Table I.

TABLE I Target composition of steels prepared (wt%) Alloy C Si Mn-Al Sn Nb si 0.05 0.3 0.64 0.03 0.3 0.9 S2 0.053 0.34 0.64 1.97 0.3 0.9 S3 0.065 0.024 0.74 2.045 0.3 0.9 S4 0.061 0.022 0.046 0. 30.9 The content of Nb and Sn were kept at the same levels as that in the Fe alloys studied in Example 2 below. The content of N was about 2 ppm, calculated according to the composition of the plain high carbon rod which was used to guarantee consistent C recovery in the melt. The content of Mn in alloy S4 and the content of Si in alloy S3 and alloy S4 were also calculated from the composition of the high carbon steel used.

All alloys were first cold rolled to approximately 50% of the starting thickness and sealed in silica tubes under a reduced pressure (0.2 atm) of argon to minimise oxidation.

Heat treatments were carried out at 850°C for 96 hours and 1150°C for 1 hour followed by air cooling. These heat treatments were chosen on the basis of assessing the stabilities of any intermetallic phases d. n- the austenite phase field and ferrite phase field respectively, longer times being required at 850°C to achieve comparable diffusion distances.

Metallography The as-cast structure of the remelted button indicated no major segregation. The overall microstructures of the steels following the different heat treatment conditions were determined using standard metallographic procedures.

SEM in-back scattered electron imaging mode was used as a first step to examine the microstructure. This technique can emphasise the compositional differences between the various phases present. Linescans of the various elements were performed using conventional electron probe microanalysis which allowed detection of very small concentrations of each element as well as giving the elemental distribution within the particles.

TEM techniques were used to characterise the precipitates present in all alloys. Thin foil specimens were prepared by twin jet electro-polishing using a solution of 300 ml ethanol, 150 ml 2-butoxyethanol and 50 ml (60%) perchloric acid. The polishing temperature was around-10°C at a voltage of 15 V. The electron microscope employed was a Philips CM 20 TEM operated at 200 kV and equipped with a energy dispersive x-ray spectrometer (EDX) system (Oxford Instruments). Whilst performing microanalysis, the specimen was held in a low background beryllium holder tilted 15° from the normal towards the detector. The X-ray count rate was optimised at about 1500-2000 counts s-lover 100 seconds. The data were acquired and analysed using a Link QX2000 system and the RTS-2/FLS program for thin foil microanalysis. This programme corrects the data for atomic number and absorption effects and takes into account the overlapping of peaks by fitting the spectra to standard profiles. K factors for calibration of the EDX spectra were taken from the virtual standards supplied with the software.

Results Microstructures of alloy S1 and alloy S2 aged at 850°C and 1150°C are shown in Figs. 1 and 2 respectively. They show that precipitates appeared in all samples and were mainly distributed along grain boundaries. For both alloys, the density of the precipitates at 1150°C (Figs. l (b) and 2 (b)) is lower than that at 850°C (Figs. l (a) and 2 (a)).

The higher Al containing samples also show more precipitates than that at lower Al compositions. As differing contrast in back scattered electron images indicates differences in composition, Figs. 2 also reveals that there is a phase which is darker than the matrix and therefore richer in the low atomic number elements.

Similar microstructures were also observed in the alloy S3 (high Al) and alloy S4 (Al free). Details of the nature of these phases will be described below.

For the lower Al containing steel (S1) aged at 850°C, EPMA line scans of C, Nb, Si and Sn (Figs. 3 (a)-3 (d)) across two types of precipitates revealed that the larger precipitate was enriched in Nb and Si while the smaller precipitates were enriched only in C and Nb. Sn was not detected in either type of precipitate. For the higher Al containing steel (S2) aged at 850°C, the line scan analyses of the same elements (Figs. 4 (a)-4 (d)) indicate that one type of bright precipitate was similar to that in the previous sample, i. e. in that it contained C and Nb.

In addition to the enrichment of Nb and Si, the other type of precipitate also contained Sn (albeit at a low level).

Fig. 5 (a)-5 (d) are the EPMA line scans of C, Nb, N and Al across the dark precipitates in the higher Al steel (S2) which revealed that these particles were enriched in N and Al. Adjacent to the large dark A1N particles are smaller precipitates enriched in Nb and C. It appears that NbC had nucleated on the surface of the A1N particles which was confirmed in the TEM observation.

After ageing at 1150°C, the C and Nb rich precipitates were found in all samples. However, A1N was still only observed in high Al content steels.

Further analysis of the microstructures of the precipitates was carried out using the selected-area diffraction (SAD) technique in the TEM. Fig. 6 (a) is a typical TEM micrograph of the C-and Nb-rich phase present in the lower Al steel aged at 1150°C. From the statistical results of more than ten different observations, the particles were found to be 0.1-1 um in diameter and ranging from near spherical to elliptical in morphology. A typical SAD pattern from this type of precipitate is shown in Figs. 6 (b), and is consistent with the NbC phase which has a f. c. c. crystal structure with a lattice parameter of a = 0.447 nm. By combining the results from SEM and EPMA, it can be concluded that C and Nb enriched phases observed in all samples are mainly NbC.

Fig. 7 (a) is a bright field TEM micrograph of the lower Al containing sample aged at 850°C. SAD patterns and EDX analyses obtained from the precipitates are consistent with the-Fe2Nb3 phase which has a fcc crystal structure with a lattice parameter of a = 1.126 nm. This type of particle normally adopts an irregular shape and has a size of about 1 to 5-FqNb3particles have also been found. A typical SAD pattern of this type of precipitate is shown in Fig. 7 (b), which is from zone axis of [Tll] h.

The microstructure of Nb, Si and Sn rich precipitates in the higher Al steel aged at 850°C is shown in Fig. 8 (a).

This micrograph is taken from the grain boundary region. A typical SAD pattern, which is consistent with the Fe2Nb- Laves phase having a hexagonal crystal structure and lattice parameters of a = 0.4830 nm and c = 0.7879 nm, is shown in Fig. 8 (b). The size of the Laves phase particles varied from 0.2 to 2 Rm It has been noticed that some Laves phase particle contained fine striations and the SAD patterns taken from them contained many streaks which were perpendicular to the fine striations suggesting a faulted structure.

Fig. 9 (a) shows a bright field micrograph of a precipitate in the higher Al steel aged at 850°C. This type of particle was found to be about 3-4 zm in diameter adopting a polygonal shape. The large A1 and N rich particles were randomly distributed in the matrix and were considered to be formed during solidification. However, these particles were found only occasionally in the TEM foils. Also this particle is normally much thicker than the adjacent matrix region within the foils and, therefore, it was difficult to produce diffraction patterns and accurate EDX analyses. To solve this problem, ion beam thinning was used to further thin the TEM foil and SAD and EDX were carried out in the thinnest areas of the samples very close to the electro-polishing hole. A typical SAD patterns from this large type of precipitate is given in Figs. 9 (b) which is consistent with the A1N with an hexagonal crystal structure and lattice parameters of a=3.11 nm and c=4.98 nm. It should also be noted (shown in Fig. 9 (a)) that there is a cluster of NbC particles nucleated on the coarse A1N particle.

TEM examinations were also carried out for alloys S3 and alloy S4 and a summary of the observations is given in Table II showing the difference between the phases observed in the four steels.

TABLE II Precipitates after ageing under given conditions Material 1 hour at 1150°C 96 hours at 850°C s si NbC NbC,-Fe2Nb3 S2 NbC, A1N NbC, A1N, Laves-Fe2Nb S3 NbC, A1N NbC, A1N, Laves-Fe2Nb S4 NbC NbC,-Fe2Nb3 It can be seen from this table that the Sn-rich Laves phase occurred in the two higher Al-containing alloys (S2 and S3) but did not appear in the lower A1 (S1) and Al-free (S4) steels in which the-metastable-FqNb3phase was found. The average of over five quantitative EDX analyses (in wt %) for Nb-rich precipitates in the four steels aged at 850°C is given in Table III (-'indicate the corresponding values are too low to be reliable).

It can be seen from Table III that the I-FeNb, phase contained the elements Fe, Nb and Si while the Fe2Nb-Laves phase contained the elements Fe, Nb, Si and, most importantly, Sn. This result is consistent with the above EPMA observations. The levels of the elements in the Laves phase and the-FqNb3phase varied due to the different matrix compositions.

TABLE III Compositions of the Nb-rich precipitates obtained by EDX in TEM (wt%) Alloys Phase Fe Nb Al Mn Si Sn s1 matrix 98.30 0.56 ~ 0.59 0.34 0.35 Fe2Nb3 29.64 66.86 ~ 0.20 3.64 ~ S2 matrix 96.71 0.23 2.02 0.47 0.36 0.21 Fe2Nb 47.07 46.73 0.33 0.46 3.94 1.46 S3 matrix 96.46 0.51 2.04 Fe2Nb 50.92 43.52 0.43 0.55 1.14 3.47 S4 matrix 98.38 0.57---0. 23 Fe2Nb3 37.26 62.41 0.34 It should be pointed out that as the concentration of A1 in the Laves phase is significant lower than in the matrix, it suggests that the Laves phase is deficient in A1 and essentially A1-free. In addition, as the level of Mn in both types of precipitates are the same order as in the corresponding matrix, the change in the precipitate compositions is mainly due to the change in the content of Fe, Nb, Si and Sn of the alloy.

Discussion Many of the particles shown in Fig. 1 and Fig. 2 have a preferential alignment, suggesting they have'nucleated on the previous grain boundaries which have a higher energy and greater atomic disorder than the grain interiors. For all alloys, the microstructure of the matrix after heat treatment at 1150°C and 850°C was bcc ferrite. However, the number density of the precipitates after heat treatment at 1150°C is lower than that at 850°C. The different rates of transformation are largely due to the fact that the solubility of the alloying elements in Fe-matrix changes with temperature.

It is well known that the solubility products of NbC and A1N increase with increasing temperature. As shown in Table II, NbC and A1N particles were found at both 850°C and 1150°C.

The solubility of Nb in Fe also decreases with decreasing temperature. Apart from the NbC and A1N particles, two types of intermetallic compound (Fe2Nb-Laves phase and-Fe2Nb3 phase) were also formed at 850°C. These will also contribute to the higher density of precipitate at lower temperature.

At both temperatures, the number density of the precipitates in the higher Al containing steel is higher than that in the lower Al composition material. As shown in Fig. 5 and Fig.

9, the A1N particles found here are much larger than the NbC particles which appear to have precipitated heterogeneously on its surface suggesting that A1N aids nucleation of the carbide precipitates by providing favourable sites. Al additions also help the formation of the stable Fe2Nb phase.

After ageing at 850°C, the intermetallic-FqNb3andFe2Nb-Laves phases were found in the low (or Al-free) and high A1- content alloys respectively. It has been reported that the -Fe2Nb3phase is not stable in the Fe-Nb system but is metastable with slow cooling rates (Bejarano et al: Z Metallkde, 82 (1991), 615 and Bejarano et al: Z Metallkde, 84 (1993), 160). The phase found in the steel S1 was found to contain about 30 Fe, 67 Nb and 4 Si (wt%), showing strong Si partition to the Fe2Nb3phase (Table III). Even in the lower Si containing steel (S4), a small amount of Si was detected in this type of particle. However, the-Fe2Nb3phase was not detected in the two higher Al containing steels.

This result suggests that it is Si not A1 that has stabilising effect on the-Fe2Nb3phase.

Thermodynamic calculations were made using the MTDATA suite of programs and thermodynamic data recommended by the SGTE and taken from the literature. Due to the limitations of the database only Fe and Nb were allowed in the Laves phase.

Thermodynamic data for the Fe2Nb3 phase was not available either, and hence it was not possible to predict the precipitation of this phase. A series of calculations were made for an alloy with an equivalent composition to alloy S3 at 850°C but with Nb contents between 0 and 2wt%. It was found that only the bcc-phase and NbC were present at all Nb levels investigated. The preferential precipitation of NbC is consistent with experiment except that it was found that with A1 additions, the Laves phase at an Nb content of 0.9 wt%. This suggests that Al accelerates the formation of the Laves phase, but without being present in it.

Al and Sn are mutually insoluble in the solid state due to an unfavourable atomic radius ratio. The enthalpies of mixing of the liquid phase (Fig. 10) calculated from the thermodynamic data are positive indicating the presence of a metastable miscibility gap at lower temperatures. On the other hand, the enthalpies of mixing of the liquid phase in the Fe-A1 system is negative. If the enthalpy of mixing curve can be taken as representative of the mutual attraction between the pairs of elements, the implication is that the Al would prefer to be in an Fe-environment (ie the ferrite matrix) rather than with Sn. Sn on the other hand would rather be with Nb, and it is those two facts which together stabilise the Laves phase.

It has been noted that Si is enriched in both the Fe2Nb- Laves phase and the-Fe2Nb3phase while Sn was only dissolved in the Fe2Nb-Laves phase. The relative stability of both types of the phases is mainly determined by the size factors. The Fe2Nb Laves phase has C14 MgZn2-type crystal structure and the ratio of the Goldschmidt atomic radii of the pure elements forming the MgZn2-type phases varies from 1.05 to 1.40. The-FqNb3intermetallic phase is a fcc of the prototype Ti2Ni and in all of this type compounds the average value of the radius ratio RA/RB for the binary phase is 1.19 and individual values do not deviate from the average by more than 6% which should be in the range The Goldschmidt atomic diameters for Fe, Nb, Sn and Si elements are calculated according to the crystal structure of elements and are listed in Table IV.

TABLE IV Lattice Spacings (a) and Goldschmidt atomic diameters (d) of the Fe, Nb, Si and Sn elements Elements Crystal a (nm) id (nm) structure (Calculated) Fe bcc (a) 0.248 0.256 Nb bcc 0.286 0.294 Si diamond 0.235 0.263 Sn diamond 0. 281 0.315 From these values, we can obtain the ratios of the Goldschmidt atomic diameters for Fe, Nb and Si which are d (Si)/d (Fe)1.03, d (Nb)/d (Si) 1.12. If Si is dissolved in the Fe2Nb Laves phase, Si atoms would replace Fe due to their similar atomic size therefore producing less strain than Si substituting for Nb. It can be seen from Table III and Fig.

11 (a) that the Fe content of the Laves phase in alloy S2 (higher in Si) is less than that in alloy S3 (lower in Si), suggesting Si additions replace Fe in the Fe2Nb-Laves phase.

Similar arguments suggest that Si additions would also replace Fe in the-FqNb3phase resulting in the lower Fe content of the-FqNb3 phase in the higher Si alloy S1 (30 wt%) than that in the lower Si alloy S4 (37 wt%), as shown in Fig. 11 (a).

The Goldschmidt atomic diameters for Fe, Nb and Sn give the atomic size ratios of d (Sn)/d (Fe) #1. 23 mdd (Sn)/d (Nb) =1. 07.

Therefore, introducing Sn in to the Fe2Nb Laves phase would result in Sn substitution for Nb. From Table III and Fig. 11 (a), it can be seen that the Nb content of the Laves phase in the lower Si alloy S3 (-44 wt % Nb) is less than that in the higher Si alloy S2 (47 wt% Nb). As Si additions replace Fe in the Laves phase, this suggests that Sn atoms would replace Nb in the Fe2Nb Laves phase rather than Fe.

In the-FtNb3phase, atomic size conditions require that Sn could only replace Nb within the-FqNM3structure. However, Sn is not detected in the-FqNb3structure. This result indicates that size factor considerations alone are inadequate to explain the phenomenon found here. Other effects such as crystal structure and atomic ordering arrangements of the phases involved should be considered as these are also importantfactors. affecting the formation and the stability of alloys phases. For example, in the hexagonal Fe2Nb Laves phase, the large Nb atoms occupy the coordination number (CN) 16 sites with 4 Nb atoms and 12 Fe atoms as neighbours. The packing of the two sets of atoms is such that only like atoms are in contact. In the fcc-FtAb3 phase, as its composition deviates from A2B stoichiometry, the atomic ordering arrangements of Nb and Fe are not clear.

In its prototype, Ti2Ni, the larger atoms (Ti) occupy two different coordinate number sites. The CN 12 sites have an equal number of Ti and Ni neighbours, and the CN 14 sites, have 10 Ti and 4 Ni neighbours. These differences might result in that Sn could replace some Nb atoms in the Fe2Nb Laves phase but not in the ii-FeNbgphase.

By comparing the compositions of the precipitates in alloys containing the same levels of Si (ie steels Sl and S2 which contain about 0.3 wt% Si and steels S3 and S4 which contain about 0.02 wt % Si) it seems that the Laves phase dissolves more Si than the corresponding-FqNb3phase (Table III and Fig. 11 (b)). It has been noted that the amounts of Sn dissolved in the Laves phase for steel S3 is higher than that in steel S2 (Fig. 11 (b)) which dissolved a higher amount of Si. This reflects the limited solubility of the third element in the Laves phase, which may be due to electronic concentration requirements as both Si and Sn are in group (IVa) of periodic table.

The concept of size factor in determining phase stability is not new but it would appear from this study that ternary additions can modify the stability of intermetallic phases in a consistent way. Hence, the solubility of Sn in the Fe2Nb Laves phase can be-understood in terms of atomic size and relative valence effect. There is also the strong suggestion that where no mutual interaction exists (eg-A1- Sn), the use of a further element, A1 in this case, can displace Sn from matrix or grain boundary sites hence offering a new strategy for dealing with Sn segregation.

Conclusions The microstructures of the precipitates in low carbon steels with Sn additions have been characterised and the effects of Al and Si on the Sn-precipitation have been discussed. The most important conclusions from this study are as follows: 1. The precipitates found in the four low carbon steels after ageing at 1150°C and 850°C comprise AlN, NbC, il-Fe2Nb3 phase and Fe2Nb-Laves phase.

2. In the lower Al and A1 free steels aged at 850°C, the metasable has been observed which dissolves Si replacing Fe. Sn remains in solution 3. In the higher A1 steels aged at 850°C, the Fe2Nb-Laves phase was found to dissolve both Si and Sn. Si additions replace Fe and Sn replaces Nb. The Laves phase has a limited solubility for the third element.

EXAMPLE 2 The formation of intermetallic compounds of tin with niobium, titanium and aluminum was investigated. The level of Nb addition was found to affect the formation of Sn-rich compounds. For small Nb-additions, no Sn containing compounds were found, whereas the higher Nb alloy exhibited a Fe2Nb-based Laves phase containing about 1% Sn.-FqNb3, associated with alumina formation was also found but did not contain tin. No Sn compounds were found with Ti additions.

Materials Two series of alloys were prepared by arc melting under an argon atmosphere. Each series was based on Fe, Sn and Al with either Nb or Ti additions. Repeated melting of the alloys was performed to ensure macroscopic homogeneity of the cast material. Care was also taken to minimise contamination by carbon, oxygen and nitrogen. Table V lists chemical analyses of the as-cast alloys.

TABLE V Chemical compositions of as-cast alloys (wt%) ALLOY Sn A1 Nb Ti Fe 1 0.055 0.44 0.12 bal.

2 0.034 0.44 0.93 bal.

3 0.048 0.44 0.086 bal.

4 0.27 0.43 0.86 bal.

All alloys were first cold rolled to approximately 50% of the starting thickness and sealed in silica tubes under a reduced pressure (0.2 atm) of argon to minimise oxidation. Heat treatment was carried out at 850°C (chosen to produce a material with a ferritic matrix but high enough to encourage the formation of any precipitates) for 96 hours during which time the material should be sufficiently homogenised, followed by air cooling.

Metallography The microstructures were studied using scanning electron microscopy (SEM), transmission electron microscopy (TEM), energy dispersive X-ray spectroscopy (EDS) and microprobe analysis.

SEM, TEM and microprobe analysis were used to determine the nature of the second-phase particles in the alloys. Specimens for SEM and microprobe analysis evaluation were mounted in non-conducting bakelite powder and polished to a 1 um finish.

Back scattered electron images were produced which emphasised the compositional differences between the various phases present. This effect is due to the strong dependence of the backscattering coefficient on the elemental atomic number.

Line scan and dot mapping of the elements were performed using conventional microprobe analysis which allowed detection of very small concentrations of each element as well as giving the elemental distribution within the particles.

TEM techniques were used to characterise the precipitates present in some samples. TEM foils were prepared by twin jet electro-polishing using a solution of 300 ml ethanol, 150 ml 2-butoxyethanol and 50 ml (60%) perchloric acid. The samples were examined using a Philips CM 20 TEM operating at 200 kV.

The precipitates were identified using selected area diffraction pattern (SADP) analysis, combined with (EDS) measurements using an ultra-thin window detector (Oxford Instruments). When performing microanalysis, the specimen was held in a low background beryllium holder tilted 15° from the normal towards the detector. The X-ray count rate was optimised at about 1500-2000 counts s'lover 100 seconds. The data were acquired and analysed using a Link QX2000 system and the RTS-2/FLS program for thin foil microanalysis. This programme corrects the data for atomic number and absorption effects and takes into account the overlapping of peaks by fitting the spectra to standard profiles.

Results Figs. 12a-d show typical SEM micrographs of the alloys studied. They show that precipitates occur in all samples and the number density of the precipitates in the high Nb and high Ti alloys (Figs. 12b and 12d) is higher than in that in the low Nb and low Ti alloys (Figs. 12a and 12c).

As the different contrast in back scattered electron images indicates differences in composition, Figs. 12a and 12b also reveal that there are two types of precipitates in both alloys. There are complex precipitates which are dark in the centre and therefore richer in the low atomic number elements and single-phase precipitates which are bright and therefore richer in-elements with an atomic number greater than that of Fe.

Microprobe line scans of 0, A1, Nb and Sn across the complex precipitates for both alloys show that the central dark particles are Al203 (Figs. 13a and 13b) and the outer part is a Nb-rich precipitate (Fig. '13c). It appears that A1203 particles act as nucleation sites for Nb-rich precipitates.

Sn was not detected in the complex precipitates (Fig. 13d).

Microprobe analysis revealed a different elemental distribution in the heavy element rich single-phase precipitates in the two alloys. For the low Nb alloy, a line scan analysis suggests that this type of precipitate is NbO and does not contain Sn (Figs. 14a-14c). However, in the high Nb alloy, the heavy element rich single-phase particles are enriched in Nb and Sn, as shown in Fig. 15a and Fig 15b.

Microprobe analysis revealed little difference between the samples with the higher and lower Ti additions. Figs. 16a-d show a line scan of Al, O, Ti and Sn across a precipitate in the high Ti alloy. The analysis shows Al203 at the centre of the precipitate (Figsl6a-b) while Ti is enriched in the outer part (Fig. 16c). Sn is not present in the precipitate (Fig. 16d). A second phase exists in the Ti containing samples and was found to be TiN as shown in Figs. 17a-b.

The preceding results are qualitative and, in order to further analyse the precipitates, TEM examinations were carried out. Fig. 18a is a bright field TEM micrograph of the high Nb sample showing the morphology of the complex precipitate. SADP and EDS analyses from the outer part of this precipitate are consistent with--the metastable q-F-e2Nb3-- phase which is cubic with a lattice parameter of a = 1.126 nm, as shown in Figs. 18b-c. A typical composition of this type of precipitate is 47.7wt% Fe, 50.8wt% Nb and 1.5wt% Al.

The composition shows an excess of Fe owing to contributions to the analysis from the matrix and the Al from the A1203 nucleus of the coprecipitate. Sn is not present in the Fe2Nb3 precipitate (Table VI).

TABLE VI: Composition of Nb-rich precipitates in Fe-0.9 Nb- 0.05 Sn-0.5 Al alloy (wt%) Precipitates Fe Nb Al Sn Fe2Nb 67.2 31.7 0.6 0.5 Fe2Nb3 47.7 50.8 1.5 matrix 98.43 0.29 1.23 0.05 Fig. 19a is a bright field TEM micrograph showing the Sn- rich precipitate morphology in the high Nb alloy. SADP analyses obtained from this precipitate are shown in Figs.

19b-c. These patterns are consistent-with the Laves phase with lattice parameters of a = 0.4830 nm and c = 0.7879 nm from zone axes of [100] and [1T0] respectively. Quantitative EDS analysis reveals that the Laves phase contains 67.2wt% Fe, 31.7wt% Nb, 0.6wt% A1 and 0.5wt% Sn (Table VI) and a typical EDS spectrum is shown in Fig. 19d (the spectrum scale has been adjusted to show the Sn peak).

Quantification of low concentrations of elements in precipitates embedded within a matrix is problematical since the exact matrix and precipitate contributions in the analysed volume are unknown. However, as the concentration of A1 in the Laves phase is lower than in the matrix it appears that the Laves phase is deficient in A1 and possibly A1-free. There would also be a slight excess of Fe in the analysis for the same reason. The results indicate that in the high Nb alloy, the Laves phase Fe2Nb dissolves some Sn.

In the Ti containing alloys, TEM observations confirmed the existence of the complex precipitate of Ti (C, O) with A1203 at the centre as indicated in Fig. 20. P was detected in the matrix (Fig. 20b) being most probably introduced as an impurity of the iron. By comparing the spectra in Figs.

20a-c it can be seen that some P also may be dissolved in the Ti (C, 0) precipitate, but not in the A1203. The most unexpected observation was the precipitation of Ti (C, O). The carbon obviously arises from contamination, probably introduced as an impurity with the Ti as C was not detected in the Nb containing alloys, which had been subject to the same experimental conditions.

Thermodynamic Calculations In parallel with the experimental work multiphase equilibrium calculations were performed in order to investigate the possibility of forming Sn compounds in the Nb and Ti alloys. This enabled the prediction of the equilibrium phases present for a variety of compositions.

As in the experimental work, two series of compositions were chosen; one comprising Nb and the other Ti. The compositions used are listed in table VII. It was also decided to include a small amount of oxygen and nitrogen in the calculations in order to simulate contamination that may have taken place during the melting process.

TABLE VII: Compositions used in multiphase thermodynamic calculations Nb containing alloy (wt%) Ti containing alloy (wt%) A1 0.5 0.5 Sn 0.055 le-5-0.3 Nb 36181 Ti 3e-5-0.9 O 0.001 0.001 N 0.002 0.002 Fe bal bal In order to be able to calculate the thermodynamic equilibrium for a given set of conditions it is necessary to use thermodynamic descriptions for all possible chemical species from the chosen components. Thermodynamic descriptions of the elements were taken from the recommendations of Dinsdale (CALPHAD, 1992,15 (4), 317- 425). Descriptions for the metallic binary systems were taken from the recommendations of the SGTE: Computer Handling and Dissemination of Data, Proc. 10th CODATA Int. Conf., Amsterdam (ed. Glaeser), 154-7; Elsvier 1987 except for the A1-Nb, Fe-Sn and Fe-Nb which were taken from the literature (Servant et al: J. Chim. Phys., 1997,94, (5), 869-888, Kumar et al CALPHAD, 1996,20, (2), 139-149, and Coelho et al in Proc. Symp. Processing and Applications of High Purity Refractory Metals and Alloys', Pittsburgh, Penn. USA, Oct 1993). No assessment of the Nb- Sn system was available at the time. Descriptions of the other compound phases such as oxides and nitrides were taken from the SGTE substance database. No descriptions for ternary solutions were available and hence all ternary interactions were considered ideal, as were binary interactions which were not available in any of the databases or in the literature. The Gibbs energy minimisation was performed using the MTDATA (Davies et al: High Temp. Sci., 1989,26,251-262) suite of programs as provided by the National Physical Laboratory, Teddington, UK.

Calculations were made using a temperature of 850°C. The stepped variable was the amount of Nb or Ti and Sn in order to mirror the experimental work. Calculations were performed on the basis of 100kg of material to provide a simple conversion to weight percent.

Fig. 24 shows the amount of phase predicted at 850°C for various amounts of Nb. As expected, the major phase present is the bcc (ferrite) phase. Also predicted in much smaller quantities are A1203 and NbN. The other phase which appears is the Laves phase Fe2Nb which is not present at low Nb concentrations, consistent with experiment. However, as no thermodynamic description is available for the Fe2 (Nbl_XSnX) phase, the Laves phase only appears as the binary compound.

Fig. 25 shows the phases in which the individual species occur. The amount of Nb dissolved in the bcc-Fe increases until saturation, then the Laves phase precipitates out. The calculation shows Sn only dissolved in the bcc-Fe, however, this is because dissolution of Sn in the Laves phase has yet to be modelled.

Figs. 26 and 27 shows similar calculation with the Ti containing mixture, but with both the amount of Ti and Sn increasing along the x-axis reflecting the experimental work. In this case, the Fe2Ti Laves phase is not predicted to appear. Ti and Sn dissolve in the matrix with only a small amount of TiN predicted to form. Again, this is consistent with experiment as no Laves phase or any other Sn containing compound was found. As with the Nb containing mixture, A1203 is predicted to form.

Discussion Both the microstructural studies and the thermodynamic calculations indicated that precipitation of elemental Sn did not occur in the alloys under study. In the Nb containing alloys, Sn was precipitated in the Fe2Nb Laves phase. This result is encouraging as it suggests that provided the Laves phase is formed then Sn could be preferentially precipitated removing its deleterious effects when segregated. Of course, the present work involved simple Fe alloys, the next stage being to include carbon and other alloying elements in the analysis.

Ti additions did not seem to have the same effect on Sn.

In fact the Sn preferred to remain in solution, even though the level of Sn was almost an order of magnitude higher than in the high-Nb alloy. Thermodynamic predictions did not suggest the formation of any FeTi phase at the alloy compositions used and none were seen experimentally.

The correlation between the thermodynamic calculations and the experimental observations is very good in the sense that the Laves phase is not predicted to appear in the low Nb containing alloy, but precipitates at higher Nb levels.

Also, no binary Sn compound was predicted to form, consistent with experiment. However, the observed solution of Sn in the Laves phase was not predicted as this was not accounted for in the calculation.

Microstructural studies revealed that complex-precipitates were present in both series of alloys after cold-rolling to 50% and heat treating at 850°C. In both cases, A1203 acts as a nucleation site for Nb-rich and Ti-rich precipitates.

Fig. 18a shows the complex precipitate in the high Nb alloy. It has been reported (J. M. Z. Bejarano, Z Metallkde., 1991,82, (8), 615-620; J. M. Z. Bejarano et al Z.

Metallkde., 1993,84, (3), 160-164; C. G. Schon, et al Intermetallics, 1996,4, (3), 211-216) that n-Fe2Nb3 is not an equilibrium phase of the Fe-Nb system but is metastable appearing during cooling. The Au203 particles are much larger than the-FqNb3particles which appear to have precipitated on its surface. The suggestion here is that during cooling, the r-Fe2Nb3particles have precipitated heterogeneously on the surface of the A1203 rather than epitaxially. This appears not to be the case for Ti (C, O) precipitation (Fig. 20) where A1203 is at the centre of the complex precipitate. The lines in the outer regions of the precipitate (the Ti (C, O)) are possibly strain lines indicating only partial coherency between the Ti (C, O) and the Al203 layers.

Conclusions Fe-Al-Nb-Sn and Fe-Al-Ti-Sn alloys with two different levels of Nb and Ti were prepared, cold rolled and heat treated at 850°C. The main findings were as follows: 1. Precipitation of Sn-containing precipitates is associated with the formation of the Fe2Nb Laves phase.

2. Precipitates which are strongly associated with the precipitation of Al203 and not containing Sn have been identified as the metastable Fe2Nb3 phase.

3. Aluminium oxide precipitates were found in all the alloys studied, acting as nucleation sites for-Fe2Nb3 or Ti (C, 0) precipitates.

4. Sn containing precipitates were not found in the Ti- containing alloys studied.

5. The experimental results are in agreement with thermodynamic predictions.

EXAMPLE 3 The basis of the approach outlined here is to study the binary interactions between the elements present within steel and to use this information to predict what may happen to the tin and copper.

Figure 28 shows the binary enthalpies of mixing for Fe with A1, Ti, Nb, Sn and Cu in the liquid phase at 1000 K calculated using the MTDATA suite of programs (Davies et al "The Application of MTDATA to the Modelling of Multicomponent Equilibria", High Temp. Sci., 26 (1990), 251- 262) and thermodynamic data taken from SGTE (% nsara et al, "Computer Handling and Dissemination of Data", Report CODATA, 1987,154-158) solution database, and ; in the case of Fe-Nb from Coelho et al"Thermodynamic Optimization of the Nb-Fe and Ta-Fe Binary systems", Klaus Schultze Symposium on Processing and Applications of High Purity Refractory Metals and Alloys, Pittsburgh, Penn. USA, Oct 1993,51-70 and for the Fe-Sn system from Kumar et al,"Thermodynamic Evaluation of Fe-Sn Phase Diagram", CALPHAD 20 (2)- (1996), 139-149.

The more negative the interaction energies the greater the attraction between the two elements. This is reflected in the binary phase diagrams for these systems. Figure 29 shows the Fe-A1 phase diagram as calculated from thermodynamic data stored in the SGTE solution database.

The diagram comprises a large bcc phase field and four compounds, namely A15Fe4, A12Fe, AlgFe2and AlFe. It is clear that this high degree of attraction seen in the liquid is also seen in the solid state.

The enthalpies of mixing for the Fe-Nb and Fe-Ti are also quite negative and both systems contain a Laves phase in the solid state. The actual positions of the phase boundaries and the size of the phase fields depends upon the competition between all of the possible phases as revealed by the minimisation of the global Gibbs energy. For example, in the Fe-Nb system (figure 30) the extent of Nb dissolution in bcc-Fe is very small in comparison with Ti dissolution in the Fe-Ti systems. This can be attributed to the relatively low Gibbs energy of mixing in the Fe-Nb bcc phase compared with the Gibbs energy of mixing in the Fe-Ti bcc phase in relation to the Gibbs energies of the Laves phases (figure 31). This is the reason why the Fe2Nb-Laves phase was found in Example 2 in mixtures of Fe-Nb-Sn-Al whereas the Fe2Ti compound was not found when Ti was substituted for Nb. In contrast, the enthalpies of mixing for liquid Fe-Sn are small and exhibit a maximum indicating immiscibility, and for Fe-Cu they are positive. Again, this is reflected in the binary phase diagrams for these two systems (figures 32 and 33). The Fe-Sn system clearly shows the miscibility gap in the liquid phase. At around 1200-1400 K a reasonably large homogeneity range of the bcc phase is observed but decreases quickly with falling temperature.

There is negligible solubility of Fe in Sn. Four compounds <BR> <BR> <BR> <BR> appear in the diagram but they have limited stability and their Gibbs energy of formation are quite low (around-2kJ to-6kJ/mol). In the Fe-Cu system (figure 33), there is negligible mutual solid solubility at low temperature. At higher temperatures, despite a maximum solubility of Cu in Fe of around 10at%, there is a miscibility gap in the fcc- phase. From this simple analysis it can be seen why Cu and Sn would tend to segregate during the steel making process.

Tin Precipitation It is proposed to make the Sn and Cu innocuous in steel by adding appropriate alloying additions to form compound phases with the two elements thus removing the grain boundary segregation. Choosing the alloying elements may be achieved by studying the thermodynamics of the binary systems.

Although an appreciation of the degree of attraction between two elements can be learned by looking at the liquid interactions, whether a binary compound is likely to form or not depends upon the Gibbs energy of the compound relative to the other competing phases. For example, whether the Fe2Ti or Fe2Nb-Laves phase forms in a Fe-rich alloy depends upon the size of the bcc phase field in each binary Fe-phase diagram and hence on the relative magnitude of the Gibbs energies of the compound and solution phases. This information is generally to hand and can allow us to predict whether a binary compound will form in a multicomponent alloy of a know composition. However, what may not be known is whether any higher order compounds or solution of third elements in the binary compounds exists.

Examination of-the enthalpies of mixing of the binary liquids comprising a third element leads to speculation of the existence of such compounds.

Figure 34 shows the enthalpy of mixing curve for the liquid phase in the Nb-Sn and the A1-Sn systems. The interaction between the Nb and Sn is very negative and again the results of this can be seen in the phase diagram with the formation of a very stable compound, namely Nb3Sn. Alternatively, the enthalpies for the Al-Sn system are very small and positive, with the suggestion of the presence of a metastable miscibility gap at lower temperatures. It can be said that Nb and Sn are quite strongly attracted to each other whereas Al and Sn are not. Thus by adding Nb to the steel it may be possible to form a Sn containing compound with the Nb. This compound is unlikely to be Nb3Sn with low levels of addition (-lwt% Nb, 0.05wt%Sn) and indeed thermodynamic calculations in Example 2 suggest that this would not be the case.

From Example 1, it seems that although Al was not associated with the Laves phase precipitate itself, it stabilised the Laves phase by encouraging the-precipitation of the Sn. This can be understood by looking at the enthalpies of mixing of the liquid phase associated with A1-Sn and Al-Fe. In the case of Al-Fe the values are large and negative indicating a high attraction between the A1 and Fe. Therefore, it would be expected for the Al to stay in the Fe matrix. However, in the case of A1 and Sn the values are small and positive indicating a repulsion. The net effect is to increase the activity of the Sn in the ferrite, enabling the Sn to force the precipitation of the stable phase through its attraction to Nb. Sn is essentially being rejected from the Fe-solution due to the presence of the Al. This effect added to the attraction between the Sn and the Nb leads to dissolution of the Sn in the Laves phase. Thus, by adding an appropriate amount of Nb to the steel and also by adding Al, Sn containing intermetallics can be formed. This can be explained by looking at the binary interactions between the different alloying elements.

Copper precipitation Although Sn is recognised to be a problem in steel it is still not known at what levels it is acceptable. Segregation of Cu however is readily found at levels of lwt%.

Experimental observations were made on materials containing Fe- 0.3wt% Cu and up to 5 wt% of A1, Si, Ti or Zr. The alloys were arc-melted, cold-rolled to 65% and annealed at 1123K for up to 2 hrs. SEM and TEM observations indicated that the Cu was no longer present as an elemental grain boundary film but in the form of precipitates comprising the microalloying element and Fe.

Figure 35 shows the binary enthalpies of mixing of the liquid phase of Cu with Fe, Ti, Al, and Si. Again, Ti, Al, and Si show negative enthalpies of mixing with Cu in the liquid phase. From the previous analysis, it would be expected that owing to these negative interactions that the alloying elements would be attracted to the copper, possibly to form intermetallic compounds.

Conclusions The simple analysis presented here works remarkably well in explaining experimental observations of the effect of adding alloying elements to steel and iron alloys containing tin and copper. The experimental observations could not be predicted by calculation of thermodynamic equilibria with the currently available thermodynamic data, and in this respect this analysis provides a useful insight into alloying behaviour. When used in the context of recycling tin-plate and incinerator scrap, it provides a possible mechanism for dealing with the problem of increasing amounts of tramp elements in steel without resorting to costly extraction processes.

The results of the analysis indicate that aluminium is a very useful alloying element for the removal of the detrimental effects of tin and copper in the steel. The implication is that aluminium waste can also be incorporated in the scrap charge along with tin-plate and incinerator scrap.

EXAMPLE 4-Hardness Study A batch of steel is prepared by melting iron and scrap steel together with alloying additions of tin, copper and aluminium giving a composition of 0.05wt% carbon, 0.3wt% tin, 2wt% copper and 4wt% aluminium. Other elements may be present, for example, manganese or silicon but at low levels such as 0.037 wt% and 0.18wt% respectively. The melting process can take place under vacuum. ; The molten material was then poured into a suitable mould and allowed to cool. The solidified ingot was then forged at 1050°C (below the melting point of copper) to encourage homogenisation of the material.

The material was then heat treated whilst encapsulated in a silica tube under a reduced pressure of argon at 1150°C for 1 hour, cooled at 950°C for 1 hour and then heated at 600°C for 16 hours. The material was found to have a higher hardness than then a similar aluminium free material heat treated in the same way.

Figure 39 illustrates a graph of Vickers Hardness for steels which underwent five different heat treatments (A, B, C, D and E). The upper line is the aluminium containing steel and the lower line is the aluminium free analogue. The increased hardness of aluminum containing steel can be seen.

Example 5-Strength Study Four steel samples were prepared with varying levels of aluminium to carry out strength measurements. The steel contained 0. 05%wt C, 0. 3%wt Sn and 2% wt Cu with an amount of aluminium shown in table VIII with the strength measurements.

Table VIII-strength measurements on four steel samples <BR> <BR> <BR> <BR> <BR> <BR> Al% wt Yield stress Ultimate Elongation% (MPa) tensile stress (MPa) 0 569 590 68 4 652 774 65 0.7 567 657 62 1.2 611 659 42. 7