What are light Alloys

advanced light alloys and composites, what elements make light alloys and what is light alloy material pdf free download
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100 Engineering Materials 2 Chapter 10 The light alloys Introduction −3 No fewer than 14 pure metals have densities 4.5 Mg m (see Table 10.1). Of these, titanium, aluminium and magnesium are in common use as structural materials. Be- ryllium is difficult to work and is toxic, but it is used in moderate quantities for heat shields and structural members in rockets. Lithium is used as an alloying element in aluminium to lower its density and save weight on airframes. Yttrium has an excellent set of properties and, although scarce, may eventually find applications in the nuclear- powered aircraft project. But the majority are unsuitable for structural use because they are chemically reactive or have low melting points. Table 10.2 shows that alloys based on aluminium, magnesium and titanium may have better stiffness/weight and strength/weight ratios than steel. Not only that; they Table 10.1 The light metals −3 Metal Density (Mg m ) T (°C) Comments m Titanium 4.50 1667 High T – excellent creep resistance. m Yttrium 4.47 1510 Good strength and ductility; scarce. Barium 3.50 729 Scandium 2.99 1538 Scarce. Aluminium 2.70 660 Strontium 2.60 770 Reactive in air/water. Caesium 1.87 28.5 Creeps/melts; very reactive in air/water. Beryllium 1.85 1287 Difficult to process; very toxic. Magnesium 1.74 649 Calcium 1.54 839 Reactive in air/water. 5 Rubidium 1.53 39 4 Sodium 0.97 98 Creep/melt; very reactive 6 Potassium 0.86 63 in air/water. 4 7 Lithium 0.53 181 There are, however, many non-structural applications for the light metals. Liquid sodium is used in large quantities for cooling nuclear reactors and in small amounts for cooling the valves of high-performance i.c. engines (it conducts heat 143 times better than water but is less dense, boils at 883°C, and is safe as long as it is kept in a sealed system.) Beryllium is used in windows for X-ray tubes. Magnesium is a catalyst for organic reactions. And the reactivity of calcium, caesium and lithium makes them useful as residual gas scavengers in vacuum systems.The light alloys 101 Table 10.2 Mechanical properties of structural light alloys 1/2 1/3 Alloy Density Young’s Yield strength E/rE /rE /r s /r Creep y −3 r(Mg m ) modulus s (MPa) temperature y E(GPa) (°C) Al alloys 2.7 71 25–600 26 3.1 1.5 9–220 150–250 Mg alloys 1.7 45 70–270 25 4.0 2.1 41–160 150–250 Ti alloys 4.5 120 170–1280 27 2.4 1.1 38–280 400–600 (Steels) (7.9) (210) (220–1600) 27 1.8 0.75 28–200 (400–600) See Chapter 25 and Fig. 25.7 for more information about these groupings. are also corrosion resistant (with titanium exceptionally so); they are non-toxic; and titanium has good creep properties. So although the light alloys were originally devel- oped for use in the aerospace industry, they are now much more widely used. The dominant use of aluminium alloys is in building and construction: panels, roofs, and frames. The second-largest consumer is the container and packaging industry; after that come transportation systems (the fastest-growing sector, with aluminium replac- ing steel and cast iron in cars and mass-transit systems); and the use of aluminium as an electrical conductor. Magnesium is lighter but more expensive. Titanium alloys are mostly used in aerospace applications where the temperatures are too high for aluminium or magnesium; but its extreme corrosion resistance makes it attractive in chemical engineering, food processing and bio-engineering. The growth in the use of these alloys is rapid: nearly 7% per year, higher than any other metals, and surpassed only by polymers. The light alloys derive their strength from solid solution hardening, age (or precip- itation) hardening, and work hardening. We now examine the principles behind each hardening mechanism, and illustrate them by drawing examples from our range of generic alloys. Solid solution hardening When other elements dissolve in a metal to form a solid solution they make the metal harder. The solute atoms differ in size, stiffness and charge from the solvent atoms. Because of this the randomly distributed solute atoms interact with dislocations and make it harder for them to move. The theory of solution hardening is rather complic- ated, but it predicts the following result for the yield strength 3/2 1/2 σ ∝ ε C , (10.1) y s where C is the solute concentration. ε is a term which represents the “mismatch” s between solute and solvent atoms. The form of this result is just what we would expect: badly matched atoms will make it harder for dislocations to move than well- matched atoms; and a large population of solute atoms will obstruct dislocations more than a sparse population.102 Engineering Materials 2 Fig. 10.1. The aluminium end of the Al–Mg phase diagram. Of the generic aluminium alloys (see Chapter 1, Table 1.4), the 5000 series derives most of its strength from solution hardening. The Al–Mg phase diagram (Fig. 10.1) shows why: at room temperature aluminium can dissolve up to 1.8 wt% magnesium at equilibrium. In practice, Al–Mg alloys can contain as much as 5.5 wt% Mg in solid solution at room temperature – a supersaturation of 5.5 − 1.8 = 3.7 wt%. In order to get this supersaturation the alloy is given the following schedule of heat treatments. (a) Hold at 450°C (“solution heat treat”) This puts the 5.5% alloy into the single phase (α) field and all the Mg will dissolve in the Al to give a random substitutional solid solution. (b) Cool moderately quickly to room temperature The phase diagram tells us that, below 275°C, the 5.5% alloy has an equilibrium struc- ture that is two-phase, α + Mg Al . If, then, we cool the alloy slowly below 275°C, Al 5 8 and Mg atoms will diffuse together to form precipitates of the intermetallic compound Mg Al . However, below 275°C, diffusion is slow and the C-curve for the precipitation 5 8 reaction is well over to the right (Fig. 10.2). So if we cool the 5.5% alloy moderately quickly we will miss the nose of the C-curve. None of the Mg will be taken out of solution as Mg Al , and we will end up with a supersaturated solid solution at room 5 8 temperature. As Table 10.3 shows, this supersaturated Mg gives a substantial increase in yield strength. Solution hardening is not confined to 5000 series aluminium alloys. The other alloy series all have elements dissolved in solid solution; and they are all solution strengthened to some degree. But most aluminium alloys owe their strength to fine precipitates of intermetallic compounds, and solution strengthening is not dominantThe light alloys 103 Fig. 10.2. Semi-schematic TTT diagram for the precipitation of Mg Al from the Al–5.5 wt% Mg 5 8 solid solution. Table 10.3 Yield strengths of 5000 series (Al–Mg) alloys Alloy wt% Mg s (MPa) y (annealed condition) 5005 0.8 40 5050 1.5 55 5 5052 2.5 90 4 5454 2.7 120 6 supersaturated 5083 4.5 145 4 7 5456 5.1 160 as it is in the 5000 series. Turning to the other light alloys, the most widely used titanium alloy (Ti–6 Al 4V) is dominated by solution hardening (Ti effectively dissolves about 7 wt% Al, and has complete solubility for V). Finally, magnesium alloys can be solution strengthened with Li, Al, Ag and Zn, which dissolve in Mg by between 2 and 5 wt%. Age (precipitation) hardening When the phase diagram for an alloy has the shape shown in Fig. 10.3 (a solid solub- ility that decreases markedly as the temperature falls), then the potential for age (or precipitation) hardening exists. The classic example is the Duralumins, or 2000 series aluminium alloys, which contain about 4% copper. The Al–Cu phase diagram tells us that, between 500°C and 580°C, the 4% Cu alloy is single phase: the Cu dissolves in the Al to give the random substitutional solid104 Engineering Materials 2 Fig. 10.3. The aluminium end of the Al–Cu phase diagram. Fig. 10.4. Room temperature microstructures in the Al + 4 wt% Cu alloy. (a) Produced by slow cooling from 550°C. (b) Produced by moderately fast cooling from 550°C. The precipitates in (a) are large and far apart. The precipitates in (b) are small and close together. solution α. Below 500°C the alloy enters the two-phase field of α + CuAl . As the 2 temperature decreases the amount of CuAl increases, and at room temperature the 2 equilibrium mixture is 93 wt% α + 7 wt% CuAl . Figure 10.4(a) shows the microstruc- 2 ture that we would get by cooling an Al–4 wt% Cu alloy slowly from 550°C to room temperature. In slow cooling the driving force for the precipitation of CuAl is small 2 and the nucleation rate is low (see Fig. 8.3). In order to accommodate the equilibriumThe light alloys 105 Fig. 10.5. TTT diagram for the precipitation of CuAl from the Al + 4 wt% Cu solid solution. Note that the 2 equilibrium solubility of Cu in Al at room temperature is only 0.1 wt% (see Fig. 10.3). The quenched solution is therefore carrying 4/0.1 = 40 times as much Cu as it wants to. amount of CuAl the few nuclei that do form grow into large precipitates of CuAl 2 2 spaced well apart. Moving dislocations find it easy to avoid the precipitates and the alloy is rather soft. If, on the other hand, we cool the alloy rather quickly, we produce a much finer structure (Fig. 10.4b). Because the driving force is large the nucleation rate is high (see Fig. 8.3). The precipitates, although small, are closely spaced: they get in the way of moving dislocations and make the alloy harder. There are limits to the precipitation hardening that can be produced by direct cooling: if the cooling rate is too high we will miss the nose of the C-curve for the precipitation reaction and will not get any precipitates at all But large increases in yield strength are possible if we age harden the alloy. To age harden our Al–4 wt% Cu alloy we use the following schedule of heat treatments. (a) Solution heat treat at 550°C. This gets all the Cu into solid solution. (b) Cool rapidly to room temperature by quenching into water or oil (“quench”). We will miss the nose of the C-curve and will end up with a highly supersaturated solid solution at room temperature (Fig. 10.5). (c) Hold at 150°C for 100 hours (“age”). As Fig. 10.5 shows, the supersaturated α will transform to the equilibrium mixture of saturated α + CuAl . But it will do so under 2 a very high driving force and will give a very fine (and very strong) structure. The C-curve nose is ≈ 150°C higher for Al–4 Cu than for Al–5.5 Mg (compare Figs 10.5 and 10.2). Diffusion is faster, and a more rapid quench is needed to miss the nose.106 Engineering Materials 2The light alloys 107 Fig. 10.6. Stages in the precipitation of CuAl . Disc-shaped GP zones (b) nucleate homogeneously from 2 supersaturated solid solution (a). The disc faces are perfectly coherent with the matrix. The disc edges are also coherent, but with a large coherency strain. (c) Some of the GP zones grow to form precipitates called q″. (The remaining GP zones dissolve and transfer Cu to the growing q″ by diffusion through the matrix.) Disc faces are perfectly coherent. Disc edges are coherent, but the mismatch of lattice parameters between the q″ and the Al matrix generates coherency strain. (d) Precipitates called q′ nucleate at matrix dislocations. The q″ precipitates all dissolve and transfer Cu to the growing q′. Disc faces are still perfectly coherent with the matrix. But disc edges are now incoherent. Neither faces nor edges show coherency strain, but for different reasons. (e) Equilibrium CuAl (q) nucleates at grain boundaries and at q′–matrix interfaces. The 2 q′ precipitates all dissolve and transfer Cu to the growing q. The CuAl is completely incoherent with the 2 matrix (see structure in Fig. 2.3). Because of this it grows as rounded rather than disc-shaped particles. Figure 10.5, as we have drawn it, is oversimplified. Because the transformation is taking place at a low temperature, where the atoms are not very mobile, it is not easy for the CuAl to separate out in one go. Instead, the transformation takes place in four 2 distinct stages. These are shown in Figs 10.6(a)–(e). The progression may appear rather involved but it is a good illustration of much of the material in the earlier chapters. More importantly, each stage of the transformation has a direct effect on the yield strength.108 Engineering Materials 2 Fig. 10.7. The yield strength of quenched Al–4 wt% Cu changes dramatically during ageing at 150°C Four separate hardening mechanisms are at work during the ageing process: (a) Solid solution hardening At the start of ageing the alloy is mostly strengthened by the 4 wt% of copper that is trapped in the supersaturated α. But when the GP zones form, almost all of the Cu is removed from solution and the solution strengthening virtually disappears (Fig. 10.7). (b) Coherency stress hardening The coherency strains around the GP zones and θ″ precipitates generate stresses that help prevent dislocation movement. The GP zones give the larger hardening effect (Fig. 10.7). (c) Precipitation hardening The precipitates can obstruct the dislocations directly. But their effectiveness is limited by two things: dislocations can either cut through the precipitates, or they can bow around them (Fig. 10.8). Resistance to cutting depends on a number of factors, of which the shearing resist- ance of the precipitate lattice is only one. In fact the cutting stress increases with ageing time (Fig. 10.7). Bowing is easier when the precipitates are far apart. During ageing the precipitate spacing increases from 10 nm to 1 µ m and beyond (Fig. 10.9). The bowing stress therefore decreases with ageing time (Fig. 10.7).The light alloys 109 Fig. 10.8. Dislocations can get past precipitates by (a) cutting or (b) bowing. Fig. 10.9. The gradual increase of particle spacing with ageing time. The four hardening mechanisms add up to give the overall variation of yield strength shown in Fig. 10.7. Peak strength is reached if the transformation is stopped at θ″. If the alloy is aged some more the strength will decrease; and the only way of recovering the strength of an overaged alloy is to solution-treat it at 550°C, quench, and start again If the alloy is not aged for long enough, then it will not reach peak strength; but this can be put right by more ageing. Although we have chosen to age our alloy at 150°C, we could, in fact, have aged it at any temperature below 180°C (see Fig. 10.10). The lower the ageing temperature, the longer the time required to get peak hardness. In practice, the ageing time should be long enough to give good control of the heat treatment operation without being too long (and expensive). Finally, Table 10.4 shows that copper is not the only alloying element that can age- harden aluminium. Magnesium and titanium can be age hardened too, but not as much as aluminium.110 Engineering Materials 2 Fig. 10.10. Detailed TTT diagram for the Al–4 wt% Cu alloy. We get peak strength by ageing to give q″. The lower the ageing temperature, the longer the ageing time. Note that GP zones do not form above 180°C: if we age above this temperature we will fail to get the peak value of yield strength. Table 10.4 Yield strengths of heat-treatable alloys Alloy series Typical composition (wt%) s (MPa) y Slowly cooled Quenched and aged 2000 Al + 4 Cu + Mg, Si, Mn 130 465 6000 Al + 0.5 Mg 0.5 Si 85 210 7000 Al + 6 Zn + Mg, Cu, Mn 300 570 Work hardening Commercially pure aluminium (1000 series) and the non-heat-treatable aluminium alloys (3000 and 5000 series) are usually work hardened. The work hardening super- imposes on any solution hardening, to give considerable extra strength (Table 10.5). Work hardening is achieved by cold rolling. The yield strength increases with strain (reduction in thickness) according to n σ = Aε , (10.2) y where A and n are constants. For aluminium alloys, n lies between 1/6 and 1/3.The light alloys 111 Table 10.5 Yield strengths of work-hardened aluminium alloys Alloy number s (MPa) y Annealed “Half hard”“Hard” 1100 35 115 145 3005 65 140 185 5456 140 300 370 Thermal stability Aluminium and magnesium melt at just over 900 K. Room temperature is 0.3 T , and m 100°C is 0.4 T . Substantial diffusion can take place in these alloys if they are used for m long periods at temperatures approaching 80–100°C. Several processes can occur to reduce the yield strength: loss of solutes from supersaturated solid solution, over- ageing of precipitates and recrystallisation of cold-worked microstructures. This lack of thermal stability has some interesting consequences. During supersonic flight frictional heating can warm the skin of an aircraft to 150°C. Because of this, Rolls-Royce had to develop a special age-hardened aluminium alloy (RR58) which would not over-age during the lifetime of the Concorde supersonic airliner. When aluminium cables are fastened to copper busbars in power circuits contact resistance heating at the junction leads to interdiffusion of Cu and Al. Massive, brittle plates of CuAl form, which can lead to joint failures; and when light alloys are welded, the 2 properties of the heat-affected zone are usually well below those of the parent metal. Background reading M. F. Ashby and D. R. H. Jones, Engineering Materials I, 2nd edition, Butterworth-Heinemann, 1996, Chapters 7 (Case study 2), 10, 12 (Case study 2), 27. Further reading I. J. Polmear, Light Alloys, 3rd edition, Arnold, 1995. R. W. K. Honeycombe, The Plastic Deformation of Metals, Arnold, 1968. D. A. Porter and K. E. Easterling, Phase Transformations in Metals and Alloys, 2nd edition, Chapman and Hall, 1992. Problems 10.1 An alloy of A1–4 weight% Cu was heated to 550°C for a few minutes and was then quenched into water. Samples of the quenched alloy were aged at 150°C for112 Engineering Materials 2 various times before being quenched again. Hardness measurements taken from the re-quenched samples gave the following data: Ageing time (h) 0 10 100 200 1000 Hardness (MPa) 650 950 1200 1150 1000 Account briefly for this behaviour. Peak hardness is obtained after 100 h at 150°C. Estimate how long it would take to get peak hardness at (a) 130°C, (b) 170°C. Hint: use Fig. 10.10. 3 Answers: (a) 10 h; (b) 10 h. 10.2 A batch of 7000 series aluminium alloy rivets for an aircraft wing was inadvert- ently over-aged. What steps can be taken to reclaim this batch of rivets? 10.3 Two pieces of work-hardened 5000 series aluminium alloy plate were butt welded together by arc welding. After the weld had cooled to room temperature, a series of hardness measurements was made on the surface of the fabrication. Sketch the variation in hardness as the position of the hardness indenter passes across the weld from one plate to the other. Account for the form of the hardness profile, and indicate its practical consequences. 10.4 One of the major uses of aluminium is for making beverage cans. The body is cold-drawn from a single slug of 3000 series non-heat treatable alloy because this has the large ductility required for the drawing operation. However, the top of the can must have a much lower ductility in order to allow the ring-pull to work (the top must tear easily). Which alloy would you select for the top from Table 10.5? Explain the reasoning behind your choice. Why are non-heat treatable alloys used for can manufacture?Steels: I – carbon steels 113 Chapter 11 Steels: I – carbon steels Introduction Iron is one of the oldest known metals. Methods of extracting and working it have been practised for thousands of years, although the large-scale production of carbon steels is a development of the ninetenth century. From these carbon steels (which still account for 90% of all steel production) a range of alloy steels has evolved: the low alloy steels (containing up to 6% of chromium, nickel, etc.); the stainless steels (con- taining, typically, 18% chromium and 8% nickel) and the tool steels (heavily alloyed with chromium, molybdenum, tungsten, vanadium and cobalt). We already know quite a bit about the transformations that take place in steels and the microstructures that they produce. In this chapter we draw these features together and go on to show how they are instrumental in determining the mechanical properties of steels. We restrict ourselves to carbon steels; alloy steels are covered in Chapter 12. Carbon is the cheapest and most effective alloying element for hardening iron. We have already seen in Chapter 1 (Table 1.1) that carbon is added to iron in quantities ranging from 0.04 to 4 wt% to make low, medium and high carbon steels, and cast iron. The mechanical properties are strongly dependent on both the carbon content and on the type of heat treatment. Steels and cast iron can therefore be used in a very wide range of applications (see Table 1.1). Microstructures produced by slow cooling (“normalising”) Carbon steels as received “off the shelf” have been worked at high temperature (usu- ally by rolling) and have then been cooled slowly to room temperature (“normalised”). The room-temperature microstructure should then be close to equilibrium and can be inferred from the Fe–C phase diagram (Fig. 11.1) which we have already come across in the Phase Diagrams course (p. 342). Table 11.1 lists the phases in the Fe–Fe C system 3 and Table 11.2 gives details of the composite eutectoid and eutectic structures that occur during slow cooling. People have sometimes been able to avoid the tedious business of extracting iron from its natural ore. When Commander Peary was exploring Greenland in 1894 he was taken by an Eskimo to a place near Cape York to see a huge, half-buried meteorite. This had provided metal for Eskimo tools and weapons for over a hundred years. Meteorites usually contain iron plus about 10% nickel: a direct delivery of low-alloy iron from the heavens.114 Engineering Materials 2 Fig. 11.1. The left-hand part of the iron–carbon phase diagram. There are five phases in the Fe–Fe C 3 system: L, d, g, a and Fe C (see Table 11.1). 3 Table 11.1 Phases in the Fe–Fe C system 3 Phase Atomic Description and comments packing Liquid d.r.p. Liquid solution of C in Fe. d b.c.c. Random interstitial solid solution of C in b.c.c. Fe. Maximum solubility of 0.08 wt% C occurs at 1492°C. Pure d Fe is the stable polymorph between 1391°C and 1536°C (see Fig. 2.1). g(also called “austenite”) f.c.c. Random interstitial solid solution of C in f.c.c. Fe. Maximum solubility of 1.7 wt% C occurs at 1130°C. Pure g Fe is the stable polymorph between 914°C and 1391°C (see Fig. 2.1). a(also called “ferrite”) b.c.c. Random interstitial solid solution of C in b.c.c. Fe. Maximum solubility of 0.035 wt% C occurs at 723°C. Pure a Fe is the stable polymorph below 914°C (see Fig. 2.1). Fe C (also called “iron Complex A hard and brittle chemical compound of Fe and C containing 3 carbide” or “cementite”) 25 atomic % (6.7 wt%) C.Steels: I – carbon steels 115 Table 11.2 Composite structures produced during the slow cooling of Fe–C alloys Name of structure Description and comments Pearlite The composite eutectoid structure of alternating plates of a and Fe C produced when 3 g containing 0.80 wt% C is cooled below 723°C (see Fig. 6.7 and Phase Diagrams p. 344). Pearlite nucleates at g grain boundaries. It occurs in low, medium and high carbon steels. It is sometimes, quite wrongly, called a phase. It is not a phase but is a mixture of the two separate phases a and Fe C in the proportions of 88.5% by weight 3 of a to 11.5% by weight of Fe C. Because grains are single crystals it is wrong to say 3 that Pearlite forms in grains: we say instead that it forms in nodules. Ledeburite The composite eutectic structure of alternating plates of g and Fe C produced when 3 liquid containing 4.3 wt% C is cooled below 1130°C. Again, not a phase Ledeburite only occurs during the solidification of cast irons, and even then the g in ledeburite will transform to a + Fe C at 723°C. 3 Fig. 11.2. Microstructures during the slow cooling of pure iron from the hot working temperature. Figures 11.2–11.6 show how the room temperature microstructure of carbon steels depends on the carbon content. The limiting case of pure iron (Fig. 11.2) is straight- forward: when γ iron cools below 914°C α grains nucleate at γ grain boundaries and the microstructure transforms to α. If we cool a steel of eutectoid composition (0.80 wt% C) below 723°C pearlite nodules nucleate at grain boundaries (Fig. 11.3) and the micro- structure transforms to pearlite. If the steel contains less than 0.80% C (a hypoeutectoid steel) then the γ starts to transform as soon as the alloy enters the α + γ field (Fig. 11.4). “Primary” α nucleates at γ grain boundaries and grows as the steel is cooled from A 3116 Engineering Materials 2 Fig. 11.3. Microstructures during the slow cooling of a eutectoid steel from the hot working temperature. As a point of detail, when pearlite is cooled to room temperature, the concentration of carbon in the a decreases slightly, following the a/a + Fe C boundary. The excess carbon reacts with iron at the a–Fe C interfaces to 3 3 form more Fe C. This “plates out” on the surfaces of the existing Fe C plates which become very slightly 3 3 thicker. The composition of Fe C is independent of temperature, of course. 3 Fig. 11.4. Microstructures during the slow cooling of a hypoeutectoid steel from the hot working temperature. A is the standard labelling for the temperature at which a first appears, and A is standard for the eutectoid 3 1 temperature. Hypoeutectoid means that the carbon content is below that of a eutectoid steel (in the same sense that hypodermic means “under the skin”).Steels: I – carbon steels 117 Fig. 11.5. Microstructures during the slow cooling of a hypereutectoid steel. A is the standard labelling for cm the temperature at which Fe C first appears. Hypereutectoid means that the carbon content is above that of a 3 eutectoid steel (in the sense that a hyperactive child has an above-normal activity). Fig. 11.6. Room temperature microstructures in slowly cooled steels of different carbon contents. (a) The proportions by weight of the different phases. (b) The proportions by weight of the different structures.118 Engineering Materials 2 to A . At A the remaining γ (which is now of eutectoid composition) transforms to 1 1 pearlite as usual. The room temperature microstructure is then made up of primary α + pearlite. If the steel contains more than 0.80% C (a hypereutectoid steel) then we get a room-temperature microstructure of primary Fe C plus pearlite instead (Fig. 11.5). 3 These structural differences are summarised in Fig. 11.6. Mechanical properties of normalised carbon steels Figure 11.7 shows how the mechanical properties of normalised carbon steels change with carbon content. Both the yield strength and tensile strength increase linearly with carbon content. This is what we would expect: the Fe C acts as a strengthening phase, 3 and the proportion of Fe C in the steel is linear in carbon concentration (Fig. 11.6a). 3 The ductility, on the other hand, falls rapidly as the carbon content goes up (Fig. 11.7) because the α –Fe C interfaces in pearlite are good at nucleating cracks. 3 Fig. 11.7. Mechanical properties of normalised carbon steels. Quenched and tempered carbon steels −1 We saw in Chapter 8 that, if we cool eutectoid γ to 500°C at about 200°C s , we will miss the nose of the C-curve. If we continue to cool below 280°C the unstable γ will begin to transform to martensite. At 220°C half the γ will have transformed to martensite. And at –50°C the steel will have become completely martensitic. Hypoeutectoid and hypereutectoid steels can be quenched to give martensite in exactly the same way (although, as Fig. 11.8 shows, their C-curves are slightly different). Figure 11.9 shows that the hardness of martensite increases rapidly with carbon content. This, again, is what we would expect. We saw in Chapter 8 that martensite is a supersaturated solid solution of C in Fe. Pure iron at room temperature would be b.c.c., but the supersaturated carbon distorts the lattice, making it tetragonalSteels: I – carbon steels 119 Fig. 11.8. TTT diagrams for (a) eutectoid, (b) hypoeutectoid and (c) hypereutectoid steels. (b) and (c) show (dashed lines) the C-curves for the formation of primary a and Fe C respectively. Note that, as the carbon 3 content increases, both M and M decrease. S F Fig. 11.9. The hardness of martensite increases with carbon content because of the increasing distortion of the lattice.

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