heat treatment of steel

29 Heat treatment

29.1 General introduct ion and cross references

This chapter is composed of two sections: the first is concerned with the heat treatment of steels and the second with age hardenable aluminium alloys. All compositions are given as wt. % unless specified otherwise. Related information may be found at the following locations:

�9 Information on alloy specifications and designations--Chapter 1. �9 Crystallographic data on some of the phases discussed here, information on relevant metallo-

graphic techniques and phase diagrams--Chapters 6. 10 and 11 respectively. �9 Diffusion data---Chapter 13. �9 Data relevant to temperature measurement by thermocouple and pyrometer techniques--

Chapters 17 and 18, respectively. �9 Mechanical property data--Chapter 22. �9 Furnace design and vacuum systems--Chapter 40.

29.2 Heat treatment of steel

29.2.1 Introduction

Heat treating is defined by the IFHTSE (International Federation for Heat Treating and Surface Engineering) as: 'a process in which the entire object, or a portion thereof, is intentionally submitted to thermal cycles and, if required, to chemical and additional physical actions, in order to achieve desired (change in the) structures and properties'. ~ Krauss has added the additional caveat that 'heat treatment for the sole purpose of hot-working is excluded from the meaning of this definition'. 2 The thermal cycles referred to in this definition are the various heat treatment steps which include: stress relieving, austenitising, normalising, annealing, quenching, and tempering. Steel is heat treated to: control the microstructure, increase the strength and toughness, release residual stresses and prevent cracking, control hardness (and softening), improve machinability and to improve mechanical properties including: yield and tensile strength, corrosion resistance and creep performance. Each step of the heat treatment process is performed for a particular purpose. Taken together, these heat treatment steps are like 'links in a chain'. -~ The acceptability of the final properties are limited by the weakest link.

In this section, an overview of the metallurgy involved in heat treatment will be provided which includes a discussion of common microstructures and phase diagram interpretation. The use of ITD (isothermal transformation diagrams), also known as TTT (time-temperature-transformation) curves, and CCD (continuous cooling diagrams), also known as CCT (cooling time-transformation) curves, to predict microstructure formation during heat treatment will be discussed. This will be followed by an overview of the different heat treatment steps including: austenitising, stress relief, normalising, annealing, quenching and tempering. Quantitative equations for estimating appropriate heat treatment temperatures and times will be given, where possible.

29.2.2 Transformations in steels

Properties such as hardness, strength, ductility and toughness are dependent on microstructure and grain size. The first step in the heat treating process is to heat the steel to its austenitising temperature.

29-1

29-2 Heat treatment

The steel is then cooled rapidly to avoid the formation of ferrite and maximise the formation of martensite, which is a relatively hard transformation product, to achieve the desired as-quenched hardness.

The most common transformation products that may be formed from austenite in quench- hardenable steels are in order of formation with decreasing cooling rate: martensite, bainite, pearlite, and pearlite/proeutectoid ferrite. Each of these microstructures provides a unique combination of properties. The standard description of austenite and its transformation products used in ASTM E7 4 and other sources will be used here (the transformation products are listed in order of their formation with increasing cooling velocity, temperatures and compositions are from Chapter 11):5

1. Austenite--A solid solution of carbon or other elements in face centred cubic gamma iron (y- Fe). It is the desired solid solution microstructure produced prior to hardening. Gamma iron is the solid nonmagnetic phase of pure iron which is stable at temperatures between 912-1 394~ 4

2. Ferrite--This is a designation commonly applied to body centred cubic alpha iron (ct-Fe) containing alloying elements in solid solution. Alpha iron is the solid phase of pure iron which is stable at temperatures below 912~ It is ferromagnetic below 768~ Increasing carbon content decreases markedly the high temperature limit of this phase at equilibrium. Fully ferritic steels are only obtained when the carbon content is very low. However, ferrite may frequently be found to nucleate on austenite grain boundaries, leaving a layer of ferrite at the prior location of these boundaries.

3a. Pearli te~A metastable microstructure formed, when local austenite areas undergo the eutectoid reaction, in alloys of iron and carbon containing greater than 0.022 percent but less than 6.67 percent carbon. The structure is an aggregate consisting of alternate lamellae of ferrite and cementite, formed on slow cooling, during the eutectoid reaction. In an alloy of given composition, pearlite may be formed isothermally at temperatures below the eutectoid temperature by quenching austenite to a desired temperature (generally above 550~ and holding for a period of time necessary for transformation to occur. The interlamellar spacing is directly proportional to the transformation temperature, that is, the higher the temperature, the greater the spacing. 4

3b. Cementite--A very hard and brittle compound of iron and carbon corresponding to the empirical formula of Fe3C. It is commonly known as iron carbide and possesses a primitive orthorhombic lattice. In 'plain carbon steels', some of the iron atoms in the cementite lattice are replaced by manganese (manganese is added to plain carbon and other steels to 'tie-up' deleterious sulphur in the compound MnS). In 'alloy steels' the iron in cementite may be partially substituted by other elements such as chromium and tungsten. Cementite will often appear in hypo-eutectoid steels as distinct lamellae (as a constituent of pearlite) or as spheroids or globules of varying size (depending on the heat-treatment employed). Cementite is in metastable equilibrium and has a tendency to decompose into iron and graphite, although the reaction rate is very slow and is not normally observed in steels. 4 The highest cementite contents are observed in white cast irons. 5

3c. LedeburitewAn intimate mixture ofaustenite and cementite in metastable equilibrium, formed on relatively rapid cooling during the eutectic reaction in alloys of iron and carbon containing greater than 2.14 percent but less than 6.67 percent carbon. Further slow cooling causes decomposition into ferrite and cementite (in the form of pearlite) as a result of the eutectoid reaction. 4 (An eutectic reaction, or equilibrium, is defined as a reversible univariant transformation in which a liquid that is stable only at a superior temperature, decomposes into two or more conjugate solid phases, for example: L = fl + y. In the case of a eutectoid reaction, the high-temperature phase is a solid, for example: a = fl + F).

4. Baini tewA metastable microstructure resulting from the transformation of austenite at temperatures between those which produce pearlite and martensite. The formation of bainite can occur on continuous (slow) cooling if the transformation rate of austenite to pearlite is much slower than that of austenite to bainite. Ordinarily, bainite may be formed isothermally by quenching austenite to the desired temperature (below the range at which pearlite forms, but above that at which martensitic transformation is possible) and holding for a specific period of time necessary for transformation to occur. If the transformation temperature is just below that at which the finest pearlite is formed, typically 350~ (660~ the bainite (upper bainite) has a feathery appearance. If the temperature is just above that at which martensite is produced, the bainite (lower bainite) is acicular, slightly resembling tempered martensite. At higher reso- lutions using a transmission electron microscope, upper bainite is observed to consist of plates of cementite in a matrix of ferrite. These discontinuous plates tend to have parallel orientation in the direction of the longer dimension of the bainite areas. Lower bainite consists of ferrite needles containing carbide platelets in parallel arrays cross-striating each needle axis at an

Table 29.1 FEATURES OF THE IRON-CARBON SYSTEM

Reprinted from reference 6, p. 4--courtesy of Marcel Dekker. Inc.

Heat treatment o f steel 29-3

Phase or mixture of phases Name

Solid solution of carbon in c~-iron Solid solution of carbon in F-iron Iron carbide (Fe3C) Eutectic mixture of carbon solid solution in v-iron with iron carbide Eutectoid mixture of carbon solid solution in cr-iron ~vith iron carbide

Ferrite Austenite Cementite Ledeburite Pearlite

angle of about 60 ~ Intermediate bainite resembles upper bainite; however, the carbides are smaller and more randomly oriented. 4

5. Martensite--A metastable phase resulting from the diffusionless athermal decomposition of austenite below a certain temperature known as the Ms temperature (martensite start temperature). This is a mechanical shear transformation in which the parent and product phases have a specific crystallographic relationship. Martensitic transformation occurs during quenching from the austenitic condition, when the cooling rate of a steel is sufficiently high, such that the pearlite and bainite transformations are suppressed. Since the transformation is diffusionless, the composition of the martensite is identical with that of the austenite from which it is formed. Therefore, martensite is a supersaturated solid solution of carbon in distorted alpha iron (ferrite) having a body centred tetragonal lattice. At low austenite carbon contents, martensite forms by dislocation motion and the result is laths with a high dislocation density, dovetailed into packets (only the latter are visible in the light microscope). At high austenite carbon contents, martensite forms by twinning and the result is midrib twinned plates. Low carbon martensite packets may appear needle-like or vermiform in cross-section?The extent of transformation depends on the martensitic temperature range (M~-Mf) of the alloy concerned, since there is a distinct temperature where martensitic transformation begins (Ms) and ends (Mr). 5

Table 29.1 provides a summary of the features of the iron-carbon system. 6

IRON-CARBON (Fe-C) PHASE DIAGRAM

The fundamental elements of heat treatment design are derived from the Fe-C equilibrium phase diagram (as modified by the alloying additions for the steel of interest), since the science of heat treatment is dependent on the formation of the desired phases and microstructures from the transformation ofaustenite. For this discussion, the iron-carbon (Fe-C) phase diagram shown in Chapter 11 will be considered. Although commonly used, the term 'iron-carbon phase diagram' is strictly incor- rect when applied to a diagram that contains cementite on the extreme right of the diagram. When cementite is involved, this diagram should be more properly designated as the iron-cementite (Fe- Cm) metastable equilibrium diagram. This distinction arises from the fact that cementite is not truly a stable equilibrium phase and degrades to iron and graphite over very long periods of time, typically much longer times than those encountered in commercial practice. The dashed lines in the Fe-C phase diagram show the metastable equilibrium between Fe3C and different phases of iron. Solid lines show the equilibrium between iron and graphite, but graphitisation rarely occurs in steel. Thus, all references in the present section to the Fe-C diagram refer to the metastable Fe-Cm system. Thus, in the present section, the extent of the different phase fields as a function of temperature and composition are those governed by the metastable equilibrium between Fe and Fe3C and not the stable equilibrium between Fe and graphite.

Note, both martensite and cementite are 'metastable' phases in the sense that they both have only the appearance of equilibrium after suitable cooling to room temperature and are capable of transformation into the equilibrium phases with a suitable thermal exposure. 5 However, the thermal exposure needed to temper martensite is of a much shorter duration than that required to transform cementite to iron plus graphite.

When referring to the Fe-C system, as with any phase diagram, it is important to properly characterise constituents, phases and microstructures. A constituent is defined as ' . . . a phase or combination of phases that occurs in a characteristic configuration in an alloy microstructure'. 4 A phase is defined as 'a physically homogenous and distinct portion of a material system'. 4 A microstructure is defined as 'the structure of a material revealed by a microscope at magnifications greater than

2 9 - 4 H e a t t r e a t m e n t

Table 29.2 RECOMMENDED MAXIMUM HOT \\ORKING TEMPERATURES FOR VARIOUS STEELS

Property of the Timken Company. used xvith permission.

Temperantre SAE No. ~ : F

1008 1232 2250 1010 1232 2250 1015 1232 2250 1040 1204 2200

1118 1232 2250 1141 1204 2200

1350 1204 2200

2317 1232 2250 2340 1204 2200

2512 1232 2250

3 115 1 232 2 250 3 135 l 204 2 200 3 140 1 204 2 200

3240 1204 2200

3310 1232 2250 3316 1232 2250 3335 1232 2250

4017 1260 2300 4032 1204 2200 4047 1204 2200 4063 1177 2 150

4130 1204 2200 4132 1204 2200 4135 1204 2200 4140 1204 2200 4142 1204 2200

4320 1204 2200 4337 1204 2200 4340 1204 2200

4422 1232 2250 4427 1232 2250

4520 1232 2250

4615 1260 2300 4620 1260 2300 4640 1204 2200

4718 1 23 "~ "~__0

Temperature

S.qE .Vo. ~ C ~ F

4820 1232 2250

5060 1177 2150

5120 1232 2250 5 140 1204 2200 5160 1177 2150

51100 1121 2050 52 100 1121 2050

6120 1232 2250 6135 1232 2250 6150 1204 2200

8617 1232 2250 8620 1232 2250 8630 1204 2200 8 640 1204 2200 8650 1204 2200

8 720 1232 2250 8 735 1204 2200 8740 1204 2200 9310 1 ~'~ ~ 0

302 1204 2200 303 1204 2200 304 1204 2200 309 1177 2 150 310 1121 2050 316 1177 2150 317 1177 2150 321 1177 2150 347 1177 2150

410 1 204 2 200 416 1 204 2 200 420 1 204 2 200 430 1 149 2 100 440A 1 149 2 100 440C 1 121 2 050 443 1 149 2 100 446 1 038 1 900

C-Mo 1 260 2 300 DM 1 260 2 300 DM-2 I 260 2 300

25x'.S The phases shown in the Fe-C diagram include: molten alloy, austenite, ferrite, cementite and graphite. Pearlite and bainite are constituents and important microstructures but they are not phases!

The Fe-C diagram shows the effect of adding carbon to iron up to 7 percent by weight. Steels are iron alloys containing up approximately 2 percent of carbon by weight; most often the total carbon content is less than 1 percent by weight. If the total carbon content is > 2 percent by weight, the alloy is classified as a cast iron. The presence of carbon stabilises austenite and expands the temperature range within which this phase is stable. Carbon goes into interstitial solid solution in both ferrite and austenite, but is much more soluble in austenite (maximum solubility 2.14 wt. % at 1 148~ than in ferrite (maximum solubility 0.022 wt. % at 727"C).

Carbon solubility in ferrite and austenite is temperature dependent, as shown in the Fe-C phase diagram. When the insertion of carbon atoms into interstices of the relevant Fe structure exceeds the

H e a t t r e a t m e n t o f s t ee l

Table 29.3 CORRELATION OF HOT STEEL TE.XIPERATURE \VITH COL()UR

29-5

Temperature

~ F : C ttot steel colour

752 400 Red: visible in the dark 885 474 Red: visible in txvilight 975 525 Red: visible in daylight

1 077 581 Red: visible in sunlight 1 292 700 Dull red 1 472 800 Turning to cherry red 1 652 900 Cherry red 1 832 1 000 Bright cherry red 2 012 1 100 Orange red 2 192 1 200 Orange yelloxv 2 372 1 300 \Vhite 2 552 1 400 Brilliant xvhite 2 732 1 500 Dazzling white 2 912 1 600 Bluish ~vhite

carbon solubility, a new primitive orthorhombic crystal structure of cementite or iron carbide (Fe~C) which is capable of greater carbon solubilisation will be created.-

The austenite phase field shown in the phase diagram is the basis for selecting hot working and heat treating temperature limits for carbon steels. Annealing, normalisation and austenitisation processes are conducted in this region to facilitate the dissolution of carbon in iron. For example, the Fe-C phase diagram shows that if the steel is cooled slowly, the structure will change from austenite to ferrite and cementite (producing a ferrite + pearlite microstructure/. With faster cooling martensite is formed. Austenite is only stable at elevated temperatures, but can be retained as an unstable phase in high-carbon martensite.

The temperature range designated as the "critical temperature range" or 'transformation range" is defined as 'those ranges of temperature within which austenite forms during heating and transforms during cooling... ' .5 The transformation ranges on heating and cooling may overlap, but never coincide exactly and are dependent on the alloy composition and the rate of temperature change. Table 29.2 provides the recommended maximum hot working temperatures for various steels, s Table 29.3 provides a table showing a correlation of surface colours and approximate temperatures.

The transformation temperature indicates the limiting temperature of a transformation range. For irons and steels, the following standard terms are applied: 5

Accm--The temperature at which the transformation from cementite to austenite is completed during heating (in a hypereutectoid steel).

Ac l - -The temperature at which austenite begins to form during heating. Ac3--The temperature at which the transformation of ferrite to austenite is completed during

heating (in a hypoeutectoid steel). Ac4--The temperature at which austenite transforms to delta ferrite during heating (in a

hypoeutectoid steel). Arcm~The temperature at which the precipitation of cementite starts during cooling (in a

hypereutectoid steel). Ar l - -The temperature at which transformation of austenite to ferrite or to ferrite and cementite is

completed during cooling. Ar3--The temperature at which austenite begins to transform to ferrite during cooling (in a

hypoeutectoid steel). Ar4--The temperature at which delta ferrite transforms to austenite during cooling (in a

hypoeutectoid steel). Ar ' - -The temperature at which transformation from austenite to pearlite begins during cooling. Ar"- -The temperature at which transformation from austenite to martensite begins during cooling.

A summary Acl, Ac3, Arl and Ar~ values for different steels is provided in Table 29.4. s

STRUCTURAL CLASSIFICATION OF STEELS

The Fe-C phase diagram depicted in Chapter 11 illustrates an important characteristic in the steel composition range which is called the "eutectoid" which refers to the composition of a solid phase

2 9 - 6 Heat treatment

Table 29.4 APPROXIMATE CRITICAL TEMPERATURES AND Ms Mf POINTS FOR CARBON AND ALLOY STEELS

Property of the Timken Company, used with permission.

Heating Cooling Quench

Acl Ac3 Ar3 Arl 3 temp. Ms SAE No. ~ ~ ~ ~ :C ~ ~ ~ ~ ~ ~ ~

Mf

~ ~

1015 743 1370 852 1565 841 1545 688 1270 1020 732 1350 846 1555 824 1515 688 1270 1030 732 1350 807 1485 796 1465 688 1270 1035 732 1350 802 1475 782 1440 688 1270 - - 1040 732 1350 793 1460 771 1420 688 1270 - - 1045 732 1350 782 1440 763 1405 688 1270 1050 727 1340 771 1420 754 1390 688 1270 - - - - 1065 816 1500 274 525 1090 885 1625 216 420

1330 718 1325 799 1470 727 1340 627 1160 - - - - 1335 713 1315 793 1460 727 1340 629 1165 843 1550 338 640 1340 727 1340 771 1420 710 1310 627 1160 - - - - 1345 718 1325 771 1420 704 1300 627 1160 - -

2317 696 1285 779 1435 685 1265 574 1065 2330 693 1280 738 1360 652 1205 488 910/ - -

566 1050 2340 696 1285 732 1350 641 1185 571 1060 788 1450 304 580 2345 685 1265 724 1335 607 1125 560 1040 - - - -

2512 699 1290 760 1400 621 1150 571 1060 - - - - 2515 682 1260 760 1400 627 1160 588 1090

3 115 735 1355 816 1500 804 1480 671 1240 - - 3120 732 1350 804 1480 785 1445 666 1230 - - - - - - 3 130 729 1345 793 1460 738 1360 660 1220 - - - - 3 140 735 1355 766 1410 691 1275 663 1225 843 1550 332 630 3141 735 1355 766 1410 704 1300 657 1215 - - 3 150 735 1355 749 1380 691 1275 657 1215 - - - - - -

3310 724 1335 782 1440 668 1235 627 1160 - - - - 3316 724 1335 785 1445 668 1235 627 1160 - -

4027 738 1360 816 1500 760 1400 666 1230 - - 4032 727 1340 816 1500 732 1350 677 1250 - - 4042 727 1340 793 1460 727 1340 654 1210 816 1500 321 610 4053 710 1310 760 1400 716 1320 649 1200 - - - - - - 4063 738 1360 754 1390 660 1220 643 1190 816 1500 229 445 4068 741 1365 757 1395 657 1215 646 1195 - - - - 4118 752 1385 816 1500 766 1410 691 1275 4 130 749 1380 802 1475 732 1350 677 1250 871 1600 377 710 4 140 749 1380 793 1460 743 1370 693 1280 816 1500 338 640 4 147 816 1500 310 590 4 150 754 1390 788 1450 699 1290 674 1245 4 160 857 1575 260 500 4320 735 1355 807 1485 721 1330 449 840/ - - - -

632 1170 4340 732 1350 774 1425 660 1220 385' 725/ 843 1550 288 550

654 1210 4342 843 1550 277 530 4615 727 1340 807 1485 760 1400 649 1200 4620 704 1300 810 1490 724 1335 660 1220 4 640 718 1325 760 1400 660 1220 468 875 843 1550 338 640

610 1130 4 6951 - - 843 1550 124 255 4718 696 1285 821 1510 766 1410 649 1200 4815 696 1285 788 1450 710 1310 460 ~ 860

599 1110 4 820 699 1290 782 1440 682 1260 441 825:

599 1110

149 79

232

204

227

288

166

254

300 175

450

400

440

55O

330

490

(continued)

Heat treatment o f steel

Table 29.4 APPROXIMATE CRITICAL TEMPERATURES AND M..kit- POINTS FOR CARBON AND ALLOY STEELS--continued

2 9 - 7

Heating Cooling Quench

Acl Ac3 Ar3 Arl 3 tenlp. Ms SAE No. ~ ~ ~ ~ ~ :F :C :F oC ~ ~ ~

Mf

~ ~

5045 738 1360 777 1430 707 1305 679 1255 - - - - 5060 743 1370 766 1410 707 1305 696 1285 - - - - - - 5 120 749 1380 829 1525 793 1460 707 1305 - - 5 140 738 1360 788 1450 729 1345 666 1230 843 1550 332 630

51100 752 1385 768 1415 716 1320 704 1300 - - 52 100 727 1340 824 1515 716 1320 688 1270 849 1560 174 345 52 100 - - - - - - 899 1650 152 305 52 100 - - - - - - 949 1740 127 260

6117 760 1400 849 1560 777 1430 688 1270 899 1650 152 305 6120 766 1410 832 1530 782 1440 704 1300 949 1740 127 260 6140 . . . . . 843 1550 327 620 6150 749 1380 788 1450 746 1375 691 1275 - -

8615 738 1360 843 1550 791 1455 685 1265 - - - - - - 8 620 732 1350 829 1525 760 1400 649 1200 - - - - - - 8630 732 1350 804 1480 727 1340 654 1210 871 1600 366 690 8640 732 1350 779 1435 691 1275 632 1170 8650 718 1325 754 1390 671 1240 646 1195 86952 . . . . 816 1500 135 275 8 720 749 1380 827 1520 760 1400 649 1200 - - - - 8740 732 1350 788 1450 704 1300 638 1180 8 750 732 1350 766 1410 685 1265 643 1190

9310 713 1315 810 1490 707 1305 443 830 . . . . 582 1080

9317 704 1300 791 1455 699 1290 427 800 - - - - 93952 . . . . . 927 1700 77 170 9442 732 1350 779 1435 693 1280 643 1190 857 1575 327 620

238

238

282

210

460

460

540

410

I Represents the case of 4 600 grades of carburising steels. 2 Represents the case of 8 600 and 9 300 grades of carburising steels respectively. 3 When two temperatures are given for At1, the higher temperature represents the pearlite reaction and the loxver temperature the

bainite reaction.

which, upon cooling, undergoes a univariant t ransformat ion into two, or more, other solid phases. 4 For a carbon steel, the eutectoid point occurs at 0.77 wt. % carbon. This is the basis o f steel c lass i f icat ion into: hypoeutectoid, eutectoid and hypereutectoid steels.

Hypoeutec to id steels are those steels with less than -~0.80 wt. % carbon (strictly 0.77 wt. % C, but a less demand ing def in i t ion is used in commerc ia l practice). Hypoeutec to id steels can, upon initial cool ing f rom the austenite single phase field, exist as two different phases, proeutectoid ferrite and austenite, each wi th different carbon contents. Upon further cooling, the remain ing austenite undergoes the eutectoid reaction to ferrite plus cementi te (i.e. pearlite) and so the micros t ructure o f these steels typically exhibits proeutectoid ferrite grains and pearlite islands.

Eutectoid steels contain 0.76 wt. % carbon (in practice, steels with 0 .75-0 .85 wt. % carbon are often classif ied as eutectoid steels). These steels form as a solid solution at any temperature in the austenitic range and all carbon is dissolved in the austenite. At the critical (eutectoid) temperature o f the i ron-cement i te system (1 340~E 727~ there is a t ransformat ion f rom austenite to lamel lar pearlite. However, i f the steel is cooled slowly to a temperature just below Arl and held for a suitable duration, spheroidal cement i te particles in ferrite are obtained, instead o f lamellar pearlite. This micros t ructure is called spheroidite and exhibits improved machinabi l i ty and formability.

Hypereutec to id steels contain -~0 .8-2 .0wt . % carbon. Upon cool ing to Arcm, proeutectoid cement i te separates f rom austenite. Below 1 340~ 727~ the remain ing austenite t ransforms to pearlite. Hence, at room temperature, the micros t ructure consists o f proeutectoid cement i te and pearlite.

29-8 Heat treatment

ISOTHERMAL AND CONTINUOUS COOLING TRANSFORMATION DIAGRAMS FOR STEELS

Microstructures that are formed upon cooling and the proportions of each are dependent on the austenitisation conditions (which influence the austenite grain size and the solutioning of alloying elements in austenite), the time and temperature cooling history of the particular alloy and composition of the alloy. The transformation products formed are illustrated typically with the use of transformation diagrams, which show the temperature-time dependence of the microstructure formation process, for the alloy being studied.

Two of the most commonly used transformation diagrams are firstly TTT (time-temperature- transformation), which is also referred to as an ITD (isothermal transformation diagram) and secondly CCT (continuous cooling transformation), which is also referred to as a CCD (continuous cooling diagram). When selected properly, either of these types of diagrams can be used to predict the microstructure and hardness of a given steel after heat treatment or they may be used to design a heat treatment process when the desired microstructure and hardness are known.

TTT diagrams (ITD)

TTT diagrams are generated by heating small samples of steel to the desired austenitising temperature and then cooling rapidly to a temperature intermediate between the austenitising and the M~ temperature. After holding at this temperature for a fixed period of time, the transformation products are then determined. At any given holding temperature, holding times up to those at which the transformation is complete need to be investigated. This is done repeatedly for a series of temperatures until a TTT diagram is constructed such as that shown in Figure 29.1. TTT diagrams can only be read along the isotherms.

The fraction transformed to ferrite, pearlite and bainite in isothermal processes can be calculated from: 9

M = 1 - e x p ( - b t n)

800

700

600

o

500

t~ 400

o_ E

300

200

100

0 1

Austenite temperature = 880~

101 E0102 1103,, 10 4 105 I 1 I _ 1 !

Seconds 1 2 4 8 15 30 6'0 t . - ~ [ . , �9 l 1 _ l l

Time ,.-----,1~ Minutes 1 2 4 6 8 16 24 Hours

Figure 29.1 Isothermal tran.s?lormation di.gram (liD). also known as a time-tentt~etzztttre- tran~/brmation (TTT) diagram, o/an zmalloved steel containing 0.5". cartoon (Source. l'erlag Stahleisen m. b.H. Dfisseldor/).

Heat treatment of steel 29-9

where M is the fraction of the phase transformed, t is the time in seconds, b = 2 x 10 -'~ and n = 3. By convention, the beginning of the transformation is defined as 1% of the parent phase transformed and the end is defined as 99% of this phase transformed.

Only martensite formation occurs without diffusion. The Hougardy equation may be used to predict the amount of martensite formed in structural steels: '~~~

M = 1 - 0.929exp [ -0 .976 x 10-2(M, - T) 1~']

where M = the amount ofmartensite, M, is the martensite start temperature and T is the temperature of interest below the Ms temperature.

The accuracy of TTT diagrams with respect to the isothermal positions on the diagram is typically accepted to be + I0~ (+20~ or +10% with respect to time.

Examples of heat treatment processes where it is only appropriate to use a TTT diagram are: isothermal annealing, austempering and martempering. These processes are illustrated schematically in Figure 29.2. 9

o o

E #.

o o

1:2.

E #.

o o

: 3

E #.

8OO 700 600 5OO 400 300 200 100

(a) 0

-~ Ac 3 _ \ \ ._, Ac 1

A Ms

M

Iogt

900 t Ac3 8OO

700 600

5OO

400

300 200

100 0

(b)

ACl . . . . . . . .

Iog t

80O 700 600 5O0 400 300 200 100

0 (c)

: M s ~ ~

:

Ac 3 Ac~

w

Iogt

Figure 29.2 /smhermalproc'esseslor whic'h onh ITDs ( TTT diagrams) may tw u.~ed (;'elnmted /i'om relbrence 9. p. 546--c'mlrte.~:v ol.Uarcel Dekker. Inc. ). (a) isofl~ermal amwaling. ( h ) au.~tempering. (c) martempering.

29-10 Heat treatment

CCT diagrams (CCD)

Steel may also be continuously cooled at different specified rates using a dilatometer allowing the proportion of transformation products formed after cooling to various temperatures intermediate between the austenitising temperature and the Ms temperature to be determined. These data can then be used to construct a CCT diagram. CCT curves correlate the temperatures for each phase transformation, the amount of transformation product obtained for a given cooling rate with time and the cooling rate necessary to obtain martensite. These correlations are obtained from CCT diagrams by using the different cooling rate curves.

The 'critical cooling rate' is the time required to avoid formation of pearlite for the particular steel being quenched. As a general rule, a quenchant must produce a cooling rate equivalent to, or faster than, that rate indicated by the 'nose' of the pearlite transformation curve, to maximise the formation of martensite.

If the temperature-time cooling curves for the quenchant and the CCT curves of the steel are plotted on the same scale, then they may be superimposed to select the steel grade which will provide the desired microstructure and hardness for a given cooling condition. 9 This assumption is limited to bars up to 100 mm in diameter quenched in oil and bars up to 150 mm quenched in water.

CCT diagrams may be constructed in various forms such as those shown in Figure 29.3. Fig- ure 29.3a is a CCT diagram for an unalloyed carbon steel (AISI 1045) which provides curves for the beginning and ending of the different phase transformations. 11 Figure 29.3b, was generated for a DIN 50CrV4 (AISI 6145) steelfl In this figure, the fraction of the transformation product formed by a cooling curve is shown on the diagram and the resulting hardness is shown on the isotherm at the bottom of the diagram.

An alternative form of a CCT diagram is shown by Figure 29.3c. 9 This curve was not generated using a dilatometer but instead cooling curves were determined experimentally at different distances from the end of a Jominy test bar. The corresponding Jominy curve is shown along with a diagram for a particular quenchant and agitation condition which permits the prediction of cross-sectional hardness for a round bar. 9"12

Another form of CCT diagram, originally developed by Atkins, 13 is illustrated in Figure 29.3d. This CCT diagram was generated by determining the cooling curves of round bars of the alloy represented in different quenchant media and then determining the corresponding transformation temperatures, microstructures and hardnesses. 9 The data represented by these curves refer only to the centre of the bar being quenched. A scale of cooling rates is provided at the bottom of the diagram. These diagrams are read along vertical lines with respect to different cooling rates. This diagram is especially useful to quickly identify the relative hardenability of different steels.

There are a number of heat treatment processes where only the use ofa CCT diagram is appropriate. These include: continuous slow cooling processes such as normalising by cooling in air, direct quenching to obtain a fully martensitic structure and continuous cooling processes resulting in mixed microstructures as illustrated in Figure 29.4. 9

The Rose-Strassburg cooling law can be used to predict cooling times and temperatures for steel samples whose cross-sectional area are not excessively large. Therefore their cooling is modelled by the following relationship: 14

T = To exp [-ca]

where To = austenitising temperature, ct -- heat transfer coefficient and t --- time. A number of points should be noted:

�9 The CCT diagram is only valid for the steel composition for which it was determined. �9 It is NOT correct to assume that the area of intersection of a cooling curve with the transformation

product is equivalent to the amount of that product that is formed. �9 Scheil has shown that transformation begins later in time for a continuous cooling process than

for an isothermal process. 15 This is consistent with TTT and CCT curve comparison. �9 Since increasing the austenitisation temperature will shift the curves to longer transformation

times, it is necessary to use CCT diagrams generated at the desired austenitising temperature.

Caution: It is becoming increasingly common to see cooling curves (temperature-time profiles) for different cooling media (quenchants) such as oil, water, air and others, superimposed on either TTT or CCT diagrams. However, superimposition of such data, especially on a TTT diagram, is not a rigorously correct practice. Various errors are introduced into such analyses due to the inherently different kinetics of cooling used to obtain TTT or CCT diagrams versus the quenchants being represented. In particular, a continuous cooling curve can be superimposed on a CCT, but not on a TTT diagram.

Heat treatment of steel

1 0 0 0

900

800

700

O ~ 600

500 Q.. E ~- 400

300

200

100

(a)

Ck45 0.44% C-0.66% Mn (SAE 1042) Composition: 0.44% C-0.66% Mn-0.22% Si-0.22% P-0.029% S-0.15% Cr-0.02% V Austenitised at 880~ (1616=F)

!

0.1 1 Seconds

~ ' ~ 1 ~ ~ ~ ' ~ ' ~ " ~ O Hardness in HRc or HV

__U2EE_RAJ 10 10 2 103 104 105

t I . I , I

1 10 100 1000 Minutes t , , I ,,~

1 10 Time ~ Hours

29-11

(b)

800

700

600

0 o

5oo

400 E

300

200

100

tB~z.[%C l%Si t/oMnlO/o P ]O/o s I%CrI~0Ni"Io/oCul %V i Methodof melting Mc Quaid-Ehn

[o.431o.411o.82[o.o41[o.o1511.22[o.04;0:14 [().11 b.s_-M-- ' 4 . . . .

Austenitising temp. = 880~C Grain size (ASTM) = 10-11 NN;'N : 30

e

101 60 102 10 3 104 105 06 Seconds 1 1 l 1 ! . 1 I

1 2 4 8 15 30 60 Minutes ( I I 1 1 1 !

1 2 4 6 8 1624 Hours 1 I I I 1 1

1 2 3 4 6 1 0

Days

F i g u r e 2 9 . 3 a , b CCDs ( CCT diagrams)/br: ( a ) an unalloyed steel. DIN Ck45 (.4 ISI 1045, reprinted from reference 1 I, p. 164---courtesy o f Marcel Dekkep. Inc. ): (b)DIN 50Crl "4 (.41SI 6145, reprinted.l)'om relerence 9, p. 539--courtesy oflMarcel Dekkel: Inc. )


T e m p e r a t u r e o C o F

871 1600

760 1400

649 1200

538 1000

427 800

316 600

204 400

83 200

(c)

. . . . . . . . . . . . . 011COSOOn 120S, . . \ ~ ,,~ Au[tenitised at 1550:F[ " - ~ \~ ' Grain size No. 7 "

~ ( ~ 4 X N 12N9.1 ~ 5 . ~ m m Ac3 = 1430 F I

1/16 : 1/8 3/16\ 3/8 1/2 3/4 I 11/2 \ I n Ao = 1350:F ] [ 1/4 q F \ \ . \ \l, I !

,.~ \ ! / ~ ,'~ "~ "~ 4 L ~ " ~ B ~ ! - ~ / ' X X X qtuiendinc~'at~ede ndidsta nces f r~ "

. . . .

�9 " ~ ' - i x ~ " I

�9 N F P B - M - A r msfol Decl A--Austen,te M-, A - - _ . . , N,~ -- \ X~ k \ F~FFerrirtlet e ' , ,

B~Bainite ~N, N " ~ "1~ "1 . . . . )~5~ ~ M--Martensite ~ , ~ . , ~ , ~

1/16 i 1/8 3/16 1/4J 3~ "1/2 \

i 1 !

2 5 10 20 50 100 200 500 1000 Cool ing t ime, sec

o 55

~ 45 or-

~ 35 t-

~ 25 T

-,<

I 1 I I I [0.41 C. 0.86 Mn. 0.26 Si. 128 Ni. 0.71 Cr

[ Que~qched ~rom 1,55~': F I �9 1 I I

In 36 40 16 0 4 8 12 16 20 24 28 32

0 6.4 12.7 19.1 25.4 31.8 38.1 44.5 50.8 57.2 63.5 mm Distance from quenched end. 1/16 In.

Figure 29.3c CCDs ( CCT diagram.s ) /or. .-11S13140. together with a Jmninv cttrve ( reprinted./i'om re/erence 9. p. 548--courte~3" orMarcel Dekke1: Inc. )

29.2.3 Hardenability

Hardenabili~ has been defined as the ability of a ferrous material to develop hardness to a given depth after being austenitised and quenched. This general definition comprises two sub-definitions, the first of which is the ability to achieve a certain hardness. ]6 The ability to achieve a certain hardness level is associated with the highest attainable hardness which depends on the carbon content of the steel and more specifically on the amount of carbon dissolved in the austenite after austenitising.

This is illustrated by considering the problem of hardening of high-strength, high-carbon steels. The higher the concentration of dissolved carbon in the austenitic phase, the greater the increase in

H e a t t r e a t m e n t o f s t ee l 29-13

Figure 29.3d CCDs ( CCT diagram.s ) /or. ('ompari.~on with the correlation OtJomin.v distance aml cooling rate /br rottnd bars ~[variotts diameter~ quenched in into oil and water( c .-t531 International. 3hlterials Park, OH. used with permission)

mechanical strength after rapid cooling and transformation of the austenite into the metastable phase, martensite. Martensitic steels typically exhibit increasing hardness and strength with increasing carbon content, but they also exhibit relatively low ductility. However, with increasing carbon concentration, martensitic transformation from austenite becomes more difficult, resulting in a greater tendency for retained austenite and correspondingly lower strength.

29-14 H e a t t r e a t m e n t

800 700

o 0 600

500 "~ 4oo " 300 E i~. 200

100 0

(a)

L Ac 3

Ac~

,og

Ac___3 800 Ac 1

o O 700 600

�9 500 g P

400 E

300 200 100

0 ,,

(b) log t

O o

E

Ac 3 .,

Ac~ 800700 " ~ . . . .

600 A ~

500

100 r 0 . . . .

(c) log t

Figure 29.4 Heat treatment processes where only CCDs ( CCT diagrams) may be used (reprinted from relerence 9, p. 544 courteo" o f Marcel Dekkel; Inc. ). (a) direct quenching to obtain a jidO" martensitic microstructure," (b) slow cooling to obtain a jerrite-pearlite microstructure; (c) continuous cooling for a mixed microstructure

The second sub-definition of hardenability refers to the hardness distribution within a cross- section from the surface to the core under specified quenching conditions. In this case, hardenability depends on the quantity of carbon which is dissolved interstitially in austenite and the amount of alloying elements dissolved substitutionally in the austenite during austenitisation. Therefore, as Figure 29.5 shows, carbon concentrations in excess of 0.6% do not yield correspondingly greater strength. Also, increasing the carbon content influences the Mr temperature relative to Ms during rapid cooling as shown in Figure 29.6.17 In this figure, it is evident that for steels with carbon contents above 0.6%, the transformation of austenite to martensite will be incomplete if the cooling process is stopped at 0~ or higher.

The depth of hardening depends on the following factors:

�9 Size and shape of the cross-section. �9 Hardenability of the material. �9 Quenching conditions.

The cross-sectional shape exhibits a significant influence on heat extraction during quenching and, therefore, on the hardening depth. Heat extraction is dependent on the surface area exposed to the quenchant. Figure 29.7 can be used to convert square and rectangular cross-sections to equivalent circular cross-sections. 18

E E z

t "

C3} c- Q)

" 0

>-

100

k nlm~/

l / 50

I I

0 0.2 0.4 0.6 0.8 Carbon content (Wt. %)

I .... !

30 25 20 Nickel content (Wt. %)

II

Heat treatment of steel 2 9 - 1 5

Figure 29.5 h!lluence of the carbon content in steel on tire yield strength ( oo 6 ) after quench hardening. }'he yield strength values were obtained.from compression tests: the additional variation of nickel content causes negligible solid-solution-hardening and was selected to obtain a constant Ms temperature lor the start o f martensitic translbrmation (used with the permission o f the Association o f h'on and Steel Engineers (.4 ISE) )

540 1

425 -

9 315

E 2 0 5 -

95 Mf

- 2 0 1 1 1 0 0.2 0.4 1.0

1000

800

LL o

600

L

Q .

E 4oo ~

200

0.6 0.8 Carbon, %

0 1.2 1.4

Figure 29.6 Influence of the carbon content in plain-carbon steels on ttre temperature of the start o f martensite formation (Ms) and the end of martensite.lbrmation (Mr)

0 175

E E 150 - (D t -

125 O c -

, - 1 0 0 - O L

N 75 E

- o 50 ~- O

O

&.

Equivalent bar diameter, in.

1 2 3 4 5 6 7 I " w 1" - 7

Round bars- - - ," ._~ ! I . . . . ,~/,- / 6

Square bars - - ~'-7'~ ;' q) , I _ ~ , / , E

~ 5 "" / / / 4 ~ / ~ / Plates

j / y - O

25' 1 o

0 - 0 0 25 50 75 100 125 150 175

Equivalent bar diameter, mm

Figure 29.7 Correlation between rectangular cross-sections and their equivalent round bar and plate sections

2 9 - 1 6 H e a t t r e a t m e n t

Table 29.5 HARDENABILITY Ea, CTORS FOR CARBON (ONTENT. (iRAIN SIZE AND SELECTED ALLOYING ELEMENTS IN STEEL

ASM International, Materials Park, OH. used xvith permission.

Carbon content (%)

Grain size no. ( G )

6 7 8 .q llq~'ing element

mm inch mill inch mm inch Mn Si Ni Cr Mo

0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45 0.50 0.55 0.60 0.65 0.70 0.75 0.80 0.85 0.90 0.95 1.00

2.0676 0.0814 1 .9050 0.0750 1 .7704 0.0697 1 . 1 6 7 1.035 2.9286 0.1153 2 .7051 0.1065 2.5273 0.0995 1 . 3 3 3 1.070 3.5890 0.1413 3 .3401 0.1315 3.0785 0.1212 1 . 5 0 0 1.105 4.1224 0.1623 3.8329 0.1509 3.5560 0.1400 1 . 6 6 7 1.140 4.6228 0.1820 4 .2621 0.1678 3.9624 0.1560 1 . 8 3 3 1.175 5.0571 0 .1991 4.6965 0.1849 4.3180 0.1700 2.000 1.210 5.4712 0.2154 5.0800 0.2000 4.6787 0.1842 2.167 1.245 5.8420 0.2300 5.4102 0.2130 5.0190 0.1976 2.333 1.280 6.1976 0.2440 5.7379 0.2259 5.3086 0.2090 2.500 1.315 6.5532 0.2580 6.0452 0.2380 5.5880 0.2200 2.667 1.350 6.934 0.273 6.375 0.251 5.867 0.231 2.833 1.385 7.214 0.284 6.655 0.262 6.121 0.241 3.000 1.420 7.493 0.295 6.934 0.273 6.375 0.251 3.167 1.455 7.772 0.306 7.188 0.283 6.604 0.260 3.333 1.490 8.026 0.316 7.442 0.293 6.858 0.270 3.500 1.525 8.280 0.326 7.696 0.303 7.061 0.278 3.667 1.560 8.534 0.336 7.925 0.312 7.290 0.287 3.833 1.595 8.788 0.346 8.153 0.321 7.518 0.296 4.000 1.630 - - - - 4.167 1.665

- - 4.333 1.700

1.018 1.036 1.055 1.073 1.091 1109 1128 1146 1164 1182 1.201 1.219 1.237 1.255 1.273 1.291 1.309 1.321 1.345 1.364

1.1080 1.2160 1.3240 1.4320 1.54 1.6480 1.7560 1.8640 1.9720 2.0800 2.1880 2.2960 2.4040 2.5120 2.62 2.7280 2.8360 2.9440 3.0520 3.1600

1.15 1.30 1.45 1.60 1.75 1.90 2.05 2.20 2.35 2.50 2.65 2.80 2.95 3.10 3.25 3.40 3.55 3.70

The effect of steel composition on hardenability may be calculated in terms of the 'ideal critical diameter', Dx, which is defined as the largest bar diameter that can be quenched to produce 50% martensite at the centre, after quenching in an ' ideal ' quench, i.e. under ' infinite ' quenching severity. The ideal quench is one that reduces surface temperature of an austenitised steel to the bath temperature instantaneously. Under these conditions, the cooling rate at the centre of the bar depends only on the thermal diffusivity of the steel.

The ideal critical diameter may be calculated from:

DI = DI Base (carbon concentration and grain size) x fMn x fsi x fcr

x f,~lo x fv x fcu x f',i x f ,

where f,~ is a multiplicative factor for the particular substitutionally dissolved alloying element. The base value DI Base and one set of alloying factors are provided in Table 29.5.1'~2~ (Note: This is not an exhaustive list of alloying factors but these are commonly encountered and they permit calculations to illustrate the effect of steel chemistry variation on hardenability.) DI values for a range of steels with differing hardenability is provided in Table 29.6.1*21

Grain size refers to the dimensions of grains or crystals in a polycrystalline metal exclusive of twinned regions and subgrains when present. Grain size is usually estimated or measured on the cross-section of an aggregate of grains. Common units are: ( 1 ) average diameter, (2) average area, (3) number of grains per linear unit, (4) number of grains per unit area and (5) number of grains per unit volume.

Grain size may be determined according to ASTM Test Method E 1 12. 22 The procedures in Test Method E 112 describe the measurement of average grain size and include the comparison procedure, the planimetric (or Jeffries) procedure, and the intercept procedures. Standard comparison charts are provided. These test methods apply chiefly to single phase grain structures but they can be applied to determine the average size of a particular type of grain structure in a multiphase or multiconstituent specimen.

In addition, the test methods provided in ASTM E 112 are used to determine the average grain size of specimens with a unimodal distribution of grain areas, diameters, or intercept lengths. These distributions are approximately log normal. These test methods do not cover methods to characterise


Table 29.6 IDEAL CRITICAL DIAMETER I DII\ALLES FOR \\.\RIOUS STEELS

ASM International, Materials Park. OH. used xvith permission.

2 9 - 1 7

Di Di Dl

Steel mm inch Steel mm inch Steel mm inch

1 045 22.9-33.0 0.9-1.3 4 ! 35 H 63.5-83.8 2.5-3.3 8 625 H 40.6-61.0 1.6-2.4 1 090 30.5-40.6 1.2-1.6 4 14(I H 78.7-119.4 3.1--4.7 8 627 H 43.2-68.6 1.7-2.7 1 320H 35.6-63.5 1.4-2.5 4317 H 43.2-61.0 1.7-2.4 8630H 53.3-71.1 2.1-2.8 1 330 H 48.3-68.6 1.9-2.7 4 320 H 45.7-66.0 1.8-2.6 8 632 H 55.9-73.7 2.2-2.9 1 335 H 50.8-71.1 2.0-2.8 4 340 H 1 i 6.8-152.4 4.6-6.0 8 635 H 61.0-86.4 2.4-3.4 1 340 H 58.4-81.3 2.3-3.2 4 620 H 35.6-55.9 1.4-2.2 8 637 H 66.0-9 !.4 2.6-3.6 2 330 H 58.4-81.3 2.3-3.2 4 620 H 38.1-55.9 1.5-2.2 8 640 H 68.6-94.0 2.7-3.7 2345 63.5-81.3 2.5-3.2 4621 H 48.3-66.0 1.9-2.6 8641H 68.6-94.0 2.7-3.7 2512H 38.1-63.5 1.5-2.5 4640H 66.0-86.4 2.6-3.4 8642H 71.1-99.1 2.8-3.9 2515H 45.7-73.7 1.8-2.9 4812H 43.2-68.6 1.7--2.7 8645H 78.7-104.1 3.1-4.1 2517H 50.8-76.2 2.0-3.0 4815H 45.7-71.1 !.8-2.8 8647H 76.2-104.1 3.0-4.1 3 120H 38.1-58.4 1.5-2.3 4817H 55.9-73.7 2.2--2.9 8650H 83.8-114.3 3.3-4.5 3 130 H 50.8-71.1 2.0-2.8 4 820 H 55.9-81.3 2.2--3.2 8 720 H 45.7-61.0 1.8-2.4 3 135 H 55.9-78.7 2.2-3.1 5 120 H 30.5-48.3 1.2-1.9 8 735 H 68.6-91.4 2.7-3.6 3 140 H 66.0-86.4 2.6-3.4 5 130 H 53.3-73.7 2.1-2.9 8 740 H 68.6-94.0 2.7-3.7 3 340 203.2-254.0 8.0-10.0 5 132 H 55.9-73.7 2.2-2.9 8 742 H 76.2-101.6 3.0-4.0 4032 H 40.6-55.9 1.6-2.2 5 135 H 55.9-73.7 2.2-2.9 8 745 H 81.3-109.2 3.2-4.3 4037H 43.2-61.0 1.7-2.4 5 140H 55.9-78.7 2.2-3.1 8747H 88.9-116.8 3.5-4.6 4042 H 43.2-61.0 1.7-2.4 5 145 H 58.4-88.9 2.3-3.5 8 750 H 96.5-124.5 3.8-4.9 4 047 H 45.7-68.6 1.8-2.7 5 150 H 63.5-94.(/ 2.5-3.7 9 260 H 50.8-83.8 2.0-3.3 4047 H 43.2-61.0 1.7-2.4 5 152 H 83.8-119.4 3.3-4.7 9261 H 66.0-94.0 2.6-3.7 4053H 53.3-73.7 2.1-2.9 5 160H 71.1-101.6 2.8-4.(I 9262H 71.1-106.7 2.8-4.2 4 063 H 55.9-88.9 2.2-3.5 6 15(I H 71.1-99.1 2.8-3.9 9 437 H 6 !.0-94.0 2.4-3.7 4 068 H 58.4-91.4 2.3-3.6 8 617 H 33.(t-58.4 1.3-2.3 9 440 H 61.0-96.5 2.4-3.8 4 130H 45.7-66.0 1.8-2.6 8620H 40.6-58.4 1.6-2.3 9442 H 71.1-106.7 2.8-4.2 3 132H 45.7-63.5 1.8-2.5 8622H 4(/.6-58.4 1.6-2.3 9445H 71.1-111.8 2.8-4.4

the nature of these distributions. Characterisation of grain size in specimens with duplex grain size distributions is described in ASTM Test Methods E 1181. Measurement of individual, very coarse grains in a fine-grained matrix is described in Test Methods E 930. These test methods deal only with determination of planar grain size. i.e. characterisation of the two-dimensional grain sections revealed by the sectioning plane. Determination of spatial grain size, i.e. measurement of the size of the three-dimensional grains in the specimen volume, is beyond the scope of these test methods.

The test methods described in ASTM E 112 are techniques performed manually using either a standard series of graded chart images for the comparison method or simple templates for the manual counting methods. Utilisation of semi-automatic digitising tablets or automatic image analysers to measure grain size is described in ASTM Test Methods E 1382. 23

The ASTM grain size number (G), referred to in Table 29.5. is a grain size designation bearing a relationship to the average intercept distance ILl at 100 diameters magnification, according to the equation: 4

G = 1 0 . 0 0 - 2 l o g 2L

The smaller the ASTM grain size, the larger the diameter of the grains. The effect of quenching conditions on the depth of hardening are not only dependent on the

quenchant being used and its physical and chemical properties but also on process parameters such as bath temperature and agitation.

29.2.4 HardenabiliW measurement

There are numerous methods to estimate steel hardenability. However, two of the most common are: Jominy curve determination and Grossmann hardenability which will be discussed here.


P r o b e ~ - "

W a t e r - - - ~

e - l]tllfl

c "

t Pea r l i t e

0 ~ �9 ~ B a i n i t e

�9 ~ _ _1_ 1, , t_ ~ i M a r t e _ ns i t e a

0 200 400 600 800

Hardness (HV)

Figure 29.8 Schematic illustration o/the Jominv end-quench test and microstructural variation with increasing distance from the quenched end. The spacing o/the pearlite lamellae increases with distance from the quenched end, i.e. along the direction o/the arrow ( reprinted /ivm re/erence 16, p. 109--courtesv of Marcel Dekker, Inc.)

JOMINY BAR END-QUENCH TEST

The most familiar and commonly used procedure for measuring steel hardenability is the Jominy bar end-quench test, as illustrated in Figure 29.8. This test has been standardised and is described in ASTM A 255, SAE J406, DIN 50191 and ISO 642. For this test, a 100mm (4 inch) long by 25 mm (1 inch) diameter round bar is austenitised to the appropriate temperature, dropped into a fixture, and one end rapidly quenched with 24~ (75~ water from a 13 mm (0.5 inch) orifice under specified conditions.]6The austenitising temperature and time is selected according to the specific steel being studied. However, most steels are heated to temperatures of 870-900~ (1 600-1 650~ The cooling velocity decreases with increasing distance from the quenched end.

After quenching, parallel flats are ground on opposite sides of the bar and hardness measurements made at 1/16 in. (1.6mm) intervals along the bar as illustrated in Figure 29.9.16 The hardness as a function of distance from the quenched end is measured and plotted and, together with measurement of the relative areas of the martensite, bainite and pearlite that is formed, it is possible to compare the hardenability of different steels using Jominy curves. As the slope of the Jominy curve increases, the ability to harden the steel (hardenability) decreases. Conversely, decreasing slopes (or increasing flatness) of the Jominy curve indicates increasing hardenability (ease of hardening).

The Jominy end-quench is used to define the hardenability of carbon steels with different alloying elements, such as chromium (Cr), manganese (Mn), or molybdenum (Mo) and having different critical cooling velocities. Jominy curves for different alloy steels are provided in Figure 29.10. These curves illustrate that the unalloyed, 0.4% carbon steel exhibits a relatively small distance for martensite (high hardness) formation. The 1% Cr and 0.2% Mn steel, however, can be hardened up to a distance of 40mm. Figure 29.10 illustrates that steel hardenability is dependent on the steel chemistry. Unalloyed steels exhibit poor hardenability. Jominy curves provide an excellent indicator of relative steel hardenability.

The Jominy test provides valid data for steels having an ideal diameter from about 25 to 150 mm (1 to 6 in.). This test can be used for DI values less than 25 mm ( 1 in.) but Vickers or microhardness tests must be used to obtain readings that are closer to the quenched end of the bar and closer together than generally possible using the standard Rockwell 'C' hardness test method. 1~

The austenitising time and temperature, extent of special carbide solution in the austenite and degree of oxidation or surface decarburisation during austenitising, care and consistency of surface flat preparation and bar positioning prior to making hardness measurements are important factors that influence test results. Therefore, all tests should be conducted in compliance with the standard being followed. 16

Using the composition of the steel, it is possible to calculate the Jominy end-quench curve for a wide range of steels with excellent correlation to experimental results. In many cases, calculation is preferred over experimental determination.

0 GI I CD

"(D

"1-

60

50

40

30

20

10

270 70 18 5.6 K/s 489"124"32.3"10" F/s } Cooling rates

1/16 4/16 8/16 16/16 Distance from quenched end, in.

0

\ \

1.0 2 . 0

1 l ! 1 l 1 1 I

25 50

Distance from quenched end

3.0 in. 1 1

75 mm


Figure 29.9 Measuring hardness on the Jominv test specimen and plotting hardenabilita" curves (reprinted from re/erence 16. p. 109--courtesy o/Marcel Dekket: Inc. )

6O

F I

"--" " - ,0.2 % Mn

O 4 0 , - - -

~ 30 '

~ U o i

20 - " ~ ~ t e e / O. 4 ~ " "

�9 " ~ , ~ ,,

�9 m - -

1~ 1o ~o ~o ~o so Distance from the quenched front side (mm)

Figure 29.10 Jominv curve comparison of the hardenahilitv qlone unalloyed and a number of other d!l]erent alloy steels

29-20 H e a t t r ea tmen t

GROSSMANN HARDENABILITY

The Grossmann method of measuring hardenability utilises a number of cylindrical steel bars with different diameters each hardened in a given quenching medium. 16 After sectioning each bar at mid- length and examining it metallographically, the bar that has 50% martensite at its centre is selected, and the diameter of this bar is designated as the critical d iameter Dcrit. Other bars with diameters smaller than Dcrit will have more martensite and correspondingly higher hardness values and bars with diameters larger than Dcrit will attain 50% martensite only up to a certain depth as shown in Figure 29.11.16 The D~it value is valid only for the quenching medium and conditions used to determine this value.

To determine the hardenability of a steel independently of the quenching medium, Grossmann introduced the term ideal critical diameter, DI, which is the diameter of a given steel bar that would produce 50% martensite at the centre when quenched in a bath of quenching intensity H = vc. Here H = ec indicates a hypothetical quenching intensity that reduces the temperature of heated steel to the bath temperature in zero time. Alternatively, excellent correlations with reported H-values are potentially achievable using cooling rates obtained by cooling curve analysis with 13, 25, 38 and 50ram (0.5, 1.0, 1.5 and 2.0 inch) Type 304 stainless steel probes. 24 (Ideal diameters for various steels are provided in Table 29.6.1~)

To identify a quenching medium and its condition, Grossmann introduced the Quenching Intensity (Severity) factor 'H' . Table 29.7 provides a summary of Grossmann H-Factors for different quench media and different quenching conditions. ~6"-~5 Although this data has been published in numerous reference texts for many years, it is of relatively limited quantitative value. One of the most obvious reasons is that quenchant agitation is not adequately defined with respect to mass flow rate, directionality and turbulence and is often unknown, yet it exhibits enormous effects on quench severity during quenching.

Figure 29.11 Determination qtcritical diameter Dcrit according to Grosxmann (reprinted fi'om re[erence 16, p. 97--courtes 3' o.f 3,1atz'el Dekket: Inc. )

Table 29.7 EFFECT OF AGITATION ON QUENCH SEVERITY AS INDICATED BY GROSSMANN QUENCH SEVERITY FACTORS

H-Factors; :~ ASM International. Materials Park. OH. used xvith permission.

Gro.s.wnann H-t:)lctor

Agitation Air Oil attd sah l l'ater Caustic soda or brine

None 0.02 0.25-0.3 0.9-1.0 2 Mild 0.30-0.35 1.0-1.1 2-2.2 Moderate - - 0.3 5-0.4 1.2-1.3 Good 0.4-0.5 1.4- I. 5 Strong - - 0.5-0.8 1.6-2.0 Violent - - 0.8-1.1 4 5


The Grossmann value 'H' is based on the Biot (Bi} number which interrelates the interfacial heat transfer coefficient (c~), thermal conductivity (k} and the radius (R) of the round bar being hardened:

Bi = c~/kR = HD

H = c~/(2k)

Since the Biot number is dimensionless, this expression means that the Grossmann value, H, is inversely proportional to the bar diameter. This method of numerically analyzing the quenching process presumes that heat transfer is a steady state, linear {Newtonian} cooling process. However, this is seldom the case and almost never the case in vaporisable quenchants such as oil, water and aqueous polymers. Therefore, a significant error exists in the basic assumption of the method.

Another difficulty is the determination of the H-value for a cross-section other than one measured experimentally. In fact, H-values depend on the cross-sectional size. Values of H do not account for specific quenching characteristics such as composition, oil viscosity, or temperature of the quenching bath. Tables of H-values do not specify the agitation rate of the quenchant, either uniformly or precisely (see Table 29.7). Therefore, although H-values are used commonly, more current and improved procedures ought to be employed when possible. For example, cooling curve analyses and the various methods of cooling curve interpretation that have been reported are all significant improvements over the use of Grossmann Hardenability factors. 24z5

29.2.5 Austenitisation

As indicated in the discussion thus far, the austenitisation process refers to the formation of austenite by heating the steel above the critical temperature for austenite formation. It is important to note that the term austenitisation means to completely transform the steel to austenite.- However, there are a number of critically important variables in the austenitisation process, two of which are heat rate and holding (soaking) time.

The heating rate is critical. There is a specific heating rate that cannot be exceeded without causing warpage or cracking since steels typically posses insufficient plasticity to accommodate increased thermal stresses in the temperature range of 250-600~C. Therefore, this is a particularly critical temperature range especially when the component has both thick and thin cross-sections. Therefore the heating rate is dependent on: '}

�9 Size and shape of the component. �9 Initial microstructure. �9 Steel composition.

For this reason, steel is often heated to the final austenitising temperature in steps of 93-204~ (200-400~ per hour.

Shapes with corners, or sharp edges, are also susceptible to cracking and this is known as the 'corner-effect'. If heating rates are excessive or non-uniform, the resulting thermal stresses may be sufficient to cause cracking.

The propensity for steel to crack is dependent on composition. For example, increasing carbon content increases the potential for cracking. The effect of composition on the potential for cracking can be modelled by calculation of the carbon equivalent (C~q) using the following equation: ̀}

Mn Cr Mo Ni V Si - 0.5 Ti W A1 Ceq = c + - g - + - T + - f + ~ + ~ - + - - T - - + T + ~ + 10

where the elements shown represent wt. % concentrations in the steel. The limits of this equation are: C < 0.9%, Mn < 1.1%, Cr < 1.8%, Mo < 0.5~ Ni _< 5.0%, V < 0.5%, Si < 1.8%, Ti < 0.5%, W < 2.0% and A1 _< 2.0%. Crack sensitMty increases with the C~q value. The following general rules were reported by Liscic: ~

Ceq < 0.4 Steel not sensitive to cracking: may be heated quickly. Ceq = 0.4---0.7 Moderate sensitMty to cracking. Cq >_ 0.7 Steel is very sensitive to cracking and should be preheated to a temperature

close to Ac~ and held until the temperature was uniform throughout to minimise thermal stresses when austenitising.

In addition to these effects, steel with high hardness and a non-uniform microstructure should be heated more slowly, due to its crack sensitMty, than a steel with low hardness and a uniform microstructure.


Table 29.8 SAE AMS 2759/IC RECOMMENDED ANNEALING. NORMALISING, AUSTENITISING TEMPERATURES AND QUENCHANTS FOR VARIOUS STEELS

Annealing Normalising A ustenitising Material temperature temperature temperature Quenching designation I (o C) ( ~ C ) ( ~ C ) medium 3

1025 885 899 871 w,p 1035 871 899 843 o,w,p 1045 857 899 829 o,w,p 10954 816 843 802 o,p 1137 788 899 843 o, ~ p 3 140 816 899 816 o,p 4037 843 899 843 o,w,p 4 130 843 899 857 o,~;p 4 135 843 899 857 o,p 4 140 843 899 843 o,p 4 150 829 871 829 o,p 4330V 857 899 871 o,p 4335V 843 899 871 o,p 4340 843 899 816 o,p 4 640 843 899 829 o,p 6150 843 899 871 o,p 8 630 843 899 857 o, ~ p 8 735 843 899 843 o,p 8740 843 899 843 o,p

1 SAE AMS 2759/IC should be consulted for detailed description of the overall heat treatment requirements for these alloys.

2 The cooling rate is not to exceed I 1 l~ to below 538:C except for 4 330 \'i 4 335V and 4 340 to below 427~ and 4 640 to below 399~

3 o = oil, w - water and p = an aqueous polymer quenchant. 4 1095 parts should be spheroidise annealed before hardening.

The austenitisation temperature for a given steel is typically specified, such as those values shown in Table 29.8. These values were selected to provide optimum hardness and grain size. As the austenitisation temperature increases, the grain size increases. This is important because the grain size affects heat treatment and subsequent performance under various working conditions. For example, increasing grain size increases the brittle to ductile transition temperature and increases the propensity for brittle fracture. Fine grained steels have greater fatigue strength than coarse grained steels. However, coarse grained steels have better machinability than fine grained steels. The Hall-Petch equation predicts the effect of grain size (d) on yield stress (Cry):26

Cry = Cro -k- Kyd -~

where or0 and Ky are material-specific constants. Increasing the austenitising temperature also: 9

�9 Increases the hardenability due to increased carbide solubilisation and increased grain size. �9 Decreases the Ms temperature. �9 Increases the (incubation) time for isothermal transformation to pearlite or bainite to begin. �9 Increases the amount of retained austenite.

For unalloyed steels, the optimum austenitisation temperature is 30--50~ above the Ac3 temperature for hypoeutectoid steels and 30-50~ above the Ac~ for hypereutectoid steels. The alloying elements in alloy steels may shift the A~ temperature either higher or lower and therefore appropriate references such as national standards must be consulted.

A 'rule of thumb' is often used to estimate the appropriate soaking time during austenitisation:

t = 6 0 + D

where t is the soaking time in minutes, and D is the maximum diameter of the component in mil- limetres. However, such rules of thumb are imprecise. Soaking times are dependent on: geometrical factors related to the furnace and the load, type of load, type of steel, thermal properties of the load, load and furnace emissivities, initial furnace and load temperatures, characteristic fan curves, and composition of the furnace atmosphere.


Figure 29.12 Aronov load characterisation diagram lbr soaking time calculation

Aronov developed a method for predicting furnace soaking times for batch loads based on 'load charactefisation'.27-28 Load characterisation diagrams are provided in Figure 29.12 and these models are based on the generalised characterisation equation for soaking time (ts):

ts "- TsbK

where Tsb is the baseline soak temperature condition taken from Figure 29.13 and K is a correction factor, the value of which depends upon the type of steel (K - 1 for low alloy steel and 0.85 for high alloy steel). 27"28

29.2.6 Annealing

A primary purpose of annealing is to soften steel to enhance its workability and machinability. However, annealing may also be performed for: 26

�9 Relief of internal stresses arising from prior processing including casting, forging, rolling, machining and welding.

�9 Improvement or restoration of ductility and toughness. �9 Improvement of machinability. �9 Grain refinement. �9 Improvement of the uniformity of the dispersion of alloying elements. �9 Achieving a specific microstructure. �9 Reduction of gaseous content within the steel.

Annealing may be an intermediate step in an overall process or it may be the final process in the heat treatment of a component. Table 29.8 provides a summary of annealing temperatures recommended for some common steel alloys.


Figure 29.13a,b Aronov soaking times./hr. (a) packed load. (h) spaced load

Annealing processes may be classified as: full annealing, process (subcritical) annealing, isothermal annealing, recrystallisation annealing, spheroidising, and normalising. Partial (intercritical) annealing is a subclass of full annealing. Figure 29.14 provides an illustrative summary of these annealing processes.

FULL ANNEALING

Full annealing of steel involves heating the steel 30-50~ above the upper critical temperature (Ac3) for hypoeutectoid steels, followed by furnace cooling through the critical temperature range, at a specified cooling rate, which is selected based on the final microstructure and hardness required. 29 Full annealing results in a relatively coarse microstructure. This process is typically used for steels

Heat t rea tment o f s teel 29-25

Figure 29.13c,d ,4 ronov soaking times/or. (c) vertical load. (d) disks

with carbon contents of 0.30-0.60%, for example, to improve machinability. Unless otherwise noted, the term 'annealing' often refers to full annealing. ~

PARTIAL (INTERCRITICALI ANNEALING

Partial annealing, as illustrated in Figure 29.14, is conducted by heating the steel to a point within the critical temperature range (Ac ~-Ac 3 ) followed by slow furnace cooling. Partial annealing, also known as 'intercritical annealing', may be performed on hypereutectoid steels to obtain a microstructure of fine pearlite and cementite instead of coarse pearlite and a network of cementite at the grain boundaries, as observed in the case of full annealing, z~ For hypereutectoid steels, this results in grain refinement which usually occurs at 10-30:C above Ac~. Partial annealing is performed to

29-26 H e a t t rea tment

improve machinability. However, steels with a Widmanst~itten or coarse ferrite/pearlite structure are unsuitable for this process. Krauss has noted that the term 'partial annealing' is an imprecise term and to be meaningful, the type of material, the time-temperature profile of the process and the degree of cold working must be specified. ~

PROCESS (SUBCRITICAL) ANNEALING

Process annealing (Figure 29.14) is performed to improve the cold-working properties of low-carbon steels (up to 0.25% carbon) or to soften high-carbon and alloy steels to facilitate shearing, turning or straightening processes. 3~ Process annealing involves heating the steel to a temperature below (typically 10-20~ below) the lower critical temperature (Acl) and is often known as 'subcritical' annealing. After heating, the steel is cooled to room temperature in still air. The process annealing temperatures for plain carbon and low alloy steels is typically limited to about 700~ to prevent partial reaustenitisation. In some cases this is limited to about 680~ for steel compositions, such as high-nickel containing steels, where the nickel further reduces the Ac~ temperature. 3~

This process can be used to temper martensitic and bainitic microstructures to produce a softened microstructure containing spheroidal carbides in ferrite. 3~ Fine pearlite is also relatively easily softened by process annealing, while coarse pearlite is too stable to be softened by this process.

Figure 29.14 Annealing-related processes in plain carbon steels

Hea t t rea tment o f s tee l 29-27

RECRYSTALLISATION ANNEALING

Prior to cold working, steel microstructures that are either spheroidised or ferritic are highly ductile. However, when steels are cold worked, they become work hardened and ductility is reduced. Work hardening resulting from cold working can be removed by a recrystallisation process which produces strain-free grains, resulting in a ductile, spheroidised microstructure. 2 Nearly all steels which are heavily cold worked undergo recrystallisation annealing, which reduces hardness and increases ductility. 26

Recrystallisation annealing is performed by heating the steel for 30 minutes to one hour, at a temperature above the recrystallisation temperature shown in Figure 29.14. 32 When heating is complete, the steel is cooled. As opposed to other annealing processes where the processing temperature is fixed, the recrystallisation annealing temperature is dependent upon composition, prior deformation, grain size and holding time. 26 Liscic reported a correlation between recrystallisation temperature (TR) and the melting temperature (Tin) of the steel in Celsius: 9

TR -- 0.4 Tm

ISOTHERMAL ANNEALING

Isothermal annealing (Figure 29.15) is conducted by heating the steel within the austenite single- phase region (i.e. above Ac3 for a hypoeutectoid steel, or above Ac~ for a eutectoid steel) for a time sufficient to complete the solutionising process, yielding a completely austenitic microstructure. The steel is then cooled rapidly at a specified rate within the pearlite transformation range indicated by the TTT diagram for the steel (less than Ac~, typically between 600-700~ until complete transformation into ferrite plus pearlite (lamellar pearlite) occurs, at which time the steel is cooled rapidly to room temperature. ~26 It should be cautioned that it is not rigorously correct to use TTT diagrams which are developed for austenitisation temperatures which do not match those for isothermal annealing. 32

Isothermal annealing is used to achieve a more homogeneous microstructure within the steel and is faster and less expensive than full annealing. It is typically performed on hypoeutectoid steels and it is usually not performed on hypereutectoid steels. 26"3~ When isothermal annealing is used in continuous production lines for small parts or for parts with thin cross-sections, it is called 'cycle annealing'.

SPHEROIDISING (SOFT ANNEALING)

Spheroidising involves the prolonged heating of steel starting from a temperature either just above, or just below the lower critical temperature (Acl) as illustrated in Figure 29.14.

A3

A1

(a)

Figure 29.15 Schematic of a heat treatment cvcle lbr isothermal annealing of. (a) h)poeutectoid steel: (b) eutectoid steel. Note: the A3 temperature shown in the figure rel&s to that of the h)poeutectoid steel


In the simplest case, the steel is heated to just below Ac~ and held for a protracted period. Since the steel remains in the ferrite plus cementite two-phase fiel& no phase transformations occur. However, the pearlite gradually spheroidises (driven by the resulting reduction in c~-Fe3C interfacial area and hence total interfacial energy).

In commercial practice, eutectoid steels are heated initially to 20-30~C above Ac~ and hypereutectoid steels are heated initially to 30-50~ above Acl. The Acl temperature can be determined from Table 29.4, obtained from the appropriate TTT diagram or calculated from:"

Ac~(~ = 7 3 9 - 22(%C) + 2(%Si) - 7(%Mn) + 14(%Cr) + 13 (%Mo)

+ 13 (%Ni)+ 20(%V)

(Hypereutectoid alloy steels may need to be heated at a higher temperature than suggested by this equation.)

Medium carbon steels may be spheroidised by heating just above or just below the Ac~ temperature. Heating is followed by furnace cooling to a temperature just below Ar~. Instead of pearlite, the resulting microstructure consists of ferrite plus fine spheroidal and/or globular cementite (with the cementite morphology depending in part upon the carbon content of the steel). The required cooling rate is given by: 9

1. For plain carbon and low-alloy steels up to 650~C ( 1200: F)---cooling rate = 20-25 K/h (furnace cooling)

2. For medium alloy steels up to 630:C (1 166-~F)---cooling rate = 15-20 K/h (furnace cooling) 3. For high alloy steels up to 600~ (1 112~ rate= 10-15 K/h (furnace cooling)

For alloy steels, the spheroidising temperature (T) may be calculated from: '~

T(~ = 705 + 20 (%Si - %Mn + %Cr - %Mo - %Ni + %W) + 100 (%V)

High carbon and alloy steels are spheroidised to enhance their machinability and ductility. ~ This process is desirable for cold-formed low-carbon and medium-carbon steels and for high-carbon steels that are premachined, prior to final machining.

DIFFUSION (HOMOGENISING) ANNEALING

Diffusion annealing (homogenising) is performed on steel ingots and castings to minimise chemical segregation. Chemical segregation defects occur as dendrites, columnar grains and chemical occlusions. The presence of these defects produces increased brittleness and reduced ductility and toughness. The homogenisation process, illustrated in Figure 29.14, is conducted by heating the steel rapidly to 1 100-1 200~ and holding for 8 to 16 hours. The steel is then furnace cooled to 800-850~ and subsequently cooled to room temperature in still air. 3~ The defects are eliminated by solute diffusion.

NORMALISING

For normalising, hypoeutectoid steels are heated to a somewhat higher temperature (40-50:C above the Ac3 ) than that used for full annealing. Hypereutectoid steels are heated above the Acre temperature as illustrated in Figure 29.14. The holding time depends on the size of the part. The minimum time is 15 minutes at temperature, with longer times being employed for larger parts, to ensure that the part is completely austenitised. Table 29.8 provides normalisation temperatures for various commonly encountered steels.

On completion of the required holding time, specified by the size of the part, the part is cooled in still air. 1.30 For a hypoeutectoid steel, the result will be a fine ferrite plus pearlite microstructure. Plain carbon and low-alloy steels should always be normalised. 3z Some alloy steels produce martensitic microstructures even with air cooling. Therefore, with such alloys, slower cooling rates are required to provide a uniform microstructure of ferrite plus pearlite.

Normalising is conducted: 3~

�9 To provide the desired microstructure and hence mechanical properties. For a given steel composition, normalised structures will be harder and stronger with lower ductility than if fully annealed.

�9 To improve hardening response by grain refinement and improved homogenisation.


�9 To improve machining characteristics, particularly for 0.15-0.40% carbon steels. �9 To eliminate carbide networks in hypereutectoid steels.

There are various equations that may be used to calculate the hardness of normalised steel. One method is by using the Bofors equation. 3z The first step in this calculation is to determine the sum of the 'carbon potentials' - Co:

Cp -" C[1 + 0.5(C -0 .20 ) ] + Si • 0.15 -+- Mn [0.125 + 0.25 (C -0 .20 ) ]

+ P [ 1 . 2 5 - 0 . 5 ( C - 0 . 2 0 ) ] + C r • •

where C is the carbon concentration in %. The ultimate tensile strength (in MPa) after normalisation is given by:

9.81 (27 + 56Cp) for hot-rolled steel

9.81 (27 + 50Cp) for forged steel

9.81 (27 + 48Cp) for cast steel

For steels which may be used in sub-zero conditions, a double normalising treatment may be specified. 32 In these cases, the steel is first heated to 50-100:C above the usual normalising temperature. This will produce greater dissolution of the alloying elements. The second normalisation step is conducted near the lower limit of the normalisation temperature range for the purpose of producing a finer grain structure.

STRESS RELIEVING

Stress relieving is used typically to remove residual stresses which have accumulated from prior manufacturing processes. Stress relief is performed by heating to a temperature below Acl (for ferritic steels) and holding at that temperature for the required time, to achieve the desired reduction in residual stresses. The steel is then cooled sufficiently slowly to avoid the formation of excessive thermal stresses. No phase transformations occur during stress relief processing. Nayar recommends heating to: 3~

�9 550-650~ for unalloyed and low-alloy steels: �9 600-700~ for hot-work and high-speed tool steels.

These temperatures are above the recrystallisation temperatures of these types of steels. Little or no stress relief occurs at temperatures <260: C and approximately 90% of the stress is relieved at 540:C. The maximum temperature for stress relief is limited to 30:C below the tempering temperature. ~"

The results of the stress relieving process are dependent on the temperature and time which are correlated through Holloman's parameter (P):~

P = T (CIIj + log t)

where T is the temperature (K), t is the time (h) and C~j is the Holloman-Jaffe constant which is calculated from:

Cttj = 2 1 . 5 3 - (5.8 x %C)

P is a measure of the 'thermal effect" of the process and that processes with the same Holloman's parameter exhibit the same effect.

Another similar commonly used expression employed in evaluating the stress relief of spring steels is the Larson-Miller equation: ~

P = T( log t + 20)/1000

Stress relieving results in a significant reduction of yield strength in addition to a decrease in the residual stresses to some 'safe" value. Typically heating and cooling during stress relieving is performed in the furnace, particularly with distortion and crack-sensitive materials. Below 300:C, faster cooling rates can be used.


29.2.7 Quenching

QUENCHANT SELECTION AND SEVERITY

Quench severity, as expressed by the Grossman H-value (or number), is the ability of a quenching medium to extract heat from a hot steel workpiece. A typical range of Grossmann H-values (numbers) for commonly used quench media is provided in Table 29.7. Although Table 29.7 is useful to obtain a relative measure of the quench severity offered by different quench media, it is difficult to apply in practice because the actual flow rates for 'moderate', 'good', 'strong' and 'violent' agitation are unknown.

Alternatively, the measurement of actual cooling rates or heat fluxes provided by a specific quenching medium does allow a quantitative metric of the quench severity provided. Some illustrative values are provided in Table 29.9. 9.33 Typically, the greater the quench severity, the greater the propensity of a given quenching medium to cause increased distortion or cracking. This usually is the result of increased thermal stresses not transformational stresses. Specific recommendations for quench media selection for use with various steel alloys is provided by standards, such as AMS 2779. Some additional general comments regarding quenchant selection include: 3435

�9 Most machined parts made from alloy steels are oil quenched to minimize distortion. �9 Most small parts or finish-ground larger parts are free-quenched. �9 Larger gears, typically those over 8 inches are fixture (die)--quenched to control distortion. �9 Smaller gears and parts such as bushings are typically plug-quenched on a splined plug typically

constructed from carburised 8620 steel. �9 Although a reduction of quench severity leads to reduced distortion, this may also be accom-

panied by undesirable microstructures such as the formation of upper bainite within carburised parts.

�9 Quench speed may be reduced by quenching in hot 149-204~ (300--400~ oil. When hot oil quenching is used for carburised steels, lower bainite, which exhibits properties somewhat similar to those of martensite, is formed.

�9 Excellent distortion control is typically obtained with austempering, quenching into a medium just above the Ms temperature and then holding until the material transforms completely to bainite. The formation of retained austenite is a significant problem with austempering processes. Retained austenite is most pronounced where Mn and Ni are major components. The best steels for austempering are plain carbon, Cr and Mo alloy steels. 34

�9 Aqueous polymer quenchants may often be used to replace quench oils but quench severity is still of primary importance.

�9 Gas or air quenching will provide the least distortion and may be used if the steel has sufficient hardenability to provide the desired properties.

�9 Low hardenability steels are quenched in brine or vigorously agitated oil. However, even with a severe quench, undesired microstructures such as ferrite, pearlite or bainite can form.

Table 29.9 COMPARISON OF TYPICAL HEAT TRANSFER RATES

Heat tran~lbr rate Quench medium ( W'm -2 K-I ) Relerence

Furnace 15 9 Still air 30 9 Compressed air 70 9 Nitrogen ( 100 kPa) 100-150 33 Salt bath or fluidised bed 350-500 33 Nitrogen (1 MPa) 400-500 33 Air-water mixture 520 9 Helium (1 MPa) 550-600 33 Helium (2 MPa) 900-1 000 33 Still oil 1 000-1 500 33 Liquid lead 1 200 9 Hydrogen (2 MPa) 1 250-1 350 33 Circulated oil 1 800-2 200 33 Hydrogen (4 MPa) 2 100-2 300 33 Circulated water 3 000-3 500 33


It is well known that the cracking propensity increases with carbon content. Therefore, the carbon content of the steel is one of the determining factors for quenchant selection. Table 29.10 summarises the mean carbon content limits for water, brine or caustic quenching of some steels. 33"35

COMPONENT SUPPORT AND LOADING

Many parts, such as ring gears, may sag and creep under their own weight when heat treated, which is an important cause of distortion. Proper support when heating is required to minimise out-of-flatness and ovality problems which may result in long grinding times, excessive stock removal, high scrap losses and loss of case depth.36To achieve adequate distortion control, custom supports or press quenching may be required. Pinion shafts are also susceptible to bending along their length if they are improperly loaded into the furnace. When this occurs, the parts must be straightened, which will add to production cost.

SURFACE CONDITION

Quench cracking may be due to various steel related problems that are only observable after the quench but the root cause of which is not the quenching process itself. Examples include: prior steel structure, stress raisers from prior machining, laps and seams, alloy inclusion defects, grinding cracks, chemical segregation and alloying element depletion. 3' In this section, three surface condition related problems that may contribute to poor distortion control and cracking will be discussed: 'tight' scale formation, decarburisation and the formation of surface seams or 'non-metallic stringers'.

Tight scale problems are encountered with forgings hardened from direct-fired gas furnaces with high-pressure burners. 35-38 The effect of tight scale on the quenching properties of two steels, 1095 carbon steel and 18-8 stainless steel, is illustrated in Figure 29.16. 2: These cooling curves were obtained by still quenching into fast oil. A scale thickness of less than 80 ~tm (0.003 in.) increased the rate of cooling of 1095 steel, as compared with the rate obtained on a specimen without scale. However, a thick scale of 130 ~tm (0.005 in.) retarded the cooling rate. A very light scale, 13 ~tm (0.0005 in.) thick also increased the cooling rate of the 18-8 steel over that obtained with the specimen without scale.

In practice, the formation of tight scale will vary in depth over the surface of the part resulting in thermal gradients due to differences in cooling rates. This problem may yield soft spots and uncontrolled distortion and is particularly a concern with nickel-containing steels. Surface oxide formation can be minimised by the use of an appropriate protective atmosphere.

Another surface related condition is decarburisation which may lead to increased distortion or cracking. 39 At a given depth within the decarburised layer, the part does not harden as completely as it would at the same depth below the surface, if there was no decarburisation. This leads to non-uniform hardness which may contribute to increased distortion and cracking because the decarburised surface transforms at a higher temperature than the core (the M~ temperature decreases with carbon content). 3s This will lead to high residual tensile stresses at the decarburised surface or a condition of unbalanced stresses and distortion. Since the surface is decarburised, it will exhibit lower hardenability than the core. This will cause the upper transformation products to form early, nucleating additional undesirable products in the core. The decarburised side will be softer than the side that did not undergo decarburising, the greater amount of martensite in the latter leading to distortion. The solution to this problem is to restore carbon into the furnace atmosphere or machine off the decarburised layer.

Table 29.10 SUGGESTED CARBON CONTENT LIMITS FOR WATER. BRINE AND CAUSTIC QUENCHING

Hardening method Shapes Max. % Carbon

Furnace hardening

Induction hardening

General usage Simple shapes Ver3, simple shapes, e.g. bars Simple shapes Complex shapes

0.30 0.35 0.40

0.50 0.33


1000

800' o o

600

{3. E a~ 400

I - -

200

1 1095

1600

k 0lHe2mV~(~Co05 in., -

i Medium scale

0.08 mm (0.003 in.) I [ !

10 20 30 40 50 Time, s

1400 I i

1200 x_

1 0 0 0

O.. 800 E

600

400

200

1 0 0 0

800 o

_~ 600

G) {3. E 400

I - -

200

0 0 0 20 30

(a) (b) Time, s

1 18-8 stainless steel

1 1

No ~cale / /

/

I ' I Light scale

0.013 mm (0.0005 in.) I I 1

10

I i - 1600

1400 I i c

- ~2oo

- 1 0 o 0

{3. - 800 E

I - - 6OO

~ - - 400

~ 2 0 0

40 50

Figure 29.16 Cooling curves illustrating the e/li'ct Olscale on cylindrical steel bars ( 13 mm dia. x 64 ram. 0.5in. dia. x 2.5 in. ) when quenched into an accelerated oil without agitation (cop.vright.4S:~l hlternational. Materials Park. OH. used with permission t. (a) A ISI 1095 steel-oil temperature = 50 ~ C ( 125 ~ F ). (h t 18-8 stainless steel-oil temperature = 25 ~ C ( 75 = F)

Table 29.11 MINIMUM RECOMMENDED MATERIAL REMOVAL FROM HOT-ROLLED STEEL PRODUCTS TO PREVENT SURFACE SEAM AND NON-METALLIC STRINGER PROBLEMS DURING HEAT TREATMENT

.Uinimmn material removal pet" side I

Condition .Von-result)hurised Resuiphurised

Turned on centres 3~ of diameter 3.8~ of diameter Centreless turned or ground 2.6~ 3.4~

I Based on bars purchased to special straightness, i.e. 3.30ram in 1.52 m IO. 13 inch in 5 feet) maximum.

Surface seams or non-metallic inclusions, which may occur in hot-rolled or cold-finished material, are defects that prevent the hot steel from welding to itself during the forging process. These defects act as stress raisers. To prevent this problem with hot rolled bars, a layer of material should be removed before heat treatment. Recommendations made earlier by Kern are provided in Table 29.11. 3s

HEATING AND ATMOSPHERE CONTROL

An important source of steel distortion and cracking during heat treatment is non-uniform heating without the appropriate protective atmosphere. For example, if steel is heated in a direct-gas-fired furnace with high moisture content, the load being heated may absorb hydrogen leading to hydrogen embrittlement and subsequent cracking which would not normally occur with a dry atmosphere.-.4o

Localised overheating is a problem for inductively heated parts. 344~ Subsequent quenching of the part leads to quench cracks at sharp corners and areas with sudden changes in cross-sectional area (stress raisers). Cracking is due to increases in residual stresses, at the stress raisers, during the quenching process. The solution to the problem is to increase the heating rate by increasing the power density of the inductor. The temperature difference across the heated zone is decreased by continuous heating or scanning of several pistons together on a single bar. 4~

For heat treating problems related to furnace design and operation, it is usually suggested that: 4~

a. The vestibules of atmosphere-hardening furnaces should be loaded and unloaded in the presence of a purge gas. Load transfer for belt and shaker hearth furnaces should only occur with thorough purging to minimise atmosphere contamination.

b. Hardening furnaces typically contain excessive loads prior to quenching. If the steel at quenching temperature is greater than 20% of the distance from discharge to charge door, this is too large a distance. Either the production rate can be increased or some of the burners can be turned off.


Table 29.12 DIMENSIONAL VARIATION IN HARDENED HIGH-CARBON STEEL WITH TIME AT AMBIENT TEMPERATURE

29-33

Change in length (~ • 103 ) after time (days)

Steel Tempering t tardncss Ope temperature ( ~ C ) ( H RC ) 7 30 90 365

1.1% C None 66 -9.0 - 18.0 -27.0 -40.0 tool steel 120 65 -0.2 -0.6 - 1.1 - 1.9 790 ~ 205 63 0.0 -0.2 -0.3 -0.7 quench 260 62 0.0 -0.2 -0.3 -0.3

1% C/Cr None 64 -1.0 -4.2 -8.2 - 11.0 840 ~ 120 65 0.3 0.5 0.7 0.6 quench 205 62 0.0 -0.1 -0.1 -0.1

260 60 0.0 -0.1 -0.1 -0.1

RETAINED AUSTENITE

Dimensional changes, which are due to the specific volume of the transformation products formed as a result of quenching, may occur either slowly or rapidly. One of the most important, with respect to residual stress variation, distortion and cracking is the formation and transformation of retained austenite. For example, the data in Table 29.12 -~s illustrates the slow conversion of retained austenite to martensite, which was still occurring days after the original quenching process, for the two steels shown. 42"43 This is a problem since dimensional control and stability is one of the primary goals of heat treatment. Therefore, microstructural determination is an essential component of any distortion control process.

QUENCHANT UNIFORMITY

Quench non-uniformity is perhaps the greatest contributor to quench cracking. Quench non- uniformity can arise from non-uniform flow fields around the part surface during the quench or non-uniform wetting of the surface. -~3 Both lead to non-uniform heat transfer during quenching. Non-uniform quenching creates large thermal gradients between the core and the surface of the part. These two contributing factors, agitation and surface wetting will be discussed here.

Poor agitation design is a major source of quench non-uniformity. The purpose of the agitation system is not only to take hot fluid away from the surface to the heat exchanger, but also to provide uniform heat removal over the entire cooling surface of all of the parts throughout the load being quenched. Even though agitation is a critically important contributor to the performance of industrial quenching practice, relatively little is known about the quality and quantity of fluid flow encountered by the parts being quenched. Recently, agitation in various commercial quenching tanks has been studied by computational fluid dynamics (CFD) and in no case was optimal and uniform flow present, without subsequent modification of the tank. 4~ Thus, identifying the sources of non-uniform fluid flow during quenching continues to be an important tool for optimising distortion control and minimising quench cracking.

The second source of non-uniform thermal gradients during quenching is related to interfacial wetting kinematics which is of particular interest for vaporisable liquid quenchants including, water, oil and aqueous polymer solutions. Most liquid vaporisable quenchants exhibit boiling temperatures between 100 and 300=(? at atmospheric pressure. When parts are quenched in these fluids, surface wetting is usually time dependent which influences the uniformity of the cooling process and the achievable hardness and potential for the formation of soft spots. This is a problem with oil contaminated aqueous polymer quenchants, sludge contaminated oils and foaming.

29.2.8 Tempering

When steel is hardened, the as-quenched martensite is not only very, hard but also brittle. Tempering, also known as 'drawing' (not to be confused with the metal forming process of the same name), is the thermal treatment of hardened (in general quenched, but as already noted some alloy steels will harden even ifnormalised) steels to obtain the desired mechanical properties which include: improved


Table 29.13 METALLURGICAL REACTIONS OCCURRING AT VARIOUS TEMPERATURE RANGES AND RELATED PHYSICAL CHANGES OF STEEL DURING TEMPERING

Stage Temperature r a n g e Metallurgical reaction E.wansion/ Contraction

1 0-200~ 32-392 ~ F

2 200-300~ 392-572~

3 230-350~ 446-662 ~ F

4 350-700~ 662-1 292~

Precipitation of e-carbide Loss of tetragonality Decomposition of retained austenite

e-carbides decompose to cementite

Precipitation of alloy carbides. Grain coarsening

Contraction

Expansion

Contraction

Expansion

toughness and ductility, lower hardness and improved dimensional stability. During tempering, as- quenched martensite is transformed into tempered martensite which is composed of highly dispersed spheroids of cementite (or other carbides) dispersed in a soft matrix of ferrite, resulting in reduced hardness and increased toughness. The objective is to allow hardness to decrease to the desired level and then to stop the carbide decomposition by cooling. The extent of the tempering effect is determined by the temperature and time of the process.

The tempering process involves heating hardened steel to some temperature below the eutectoid temperature for the purposes of decreasing hardness and increasing toughness. (Tempering is performed as soon as possible after the steel has cooled to between 50-75~ and room temperature, to reduce the potential of cracking. If a tool steel cannot be tempered immediately after quenching, it is recommended that it be held at 50-100~ in an oven until it can be tempered. 32) In general, the tempering process is divided into four stages, which are summarised in Table 29.13. The changes occurring during tempering include: 45

1. transformation of the martensite, 2. transformation of retained austenite, 3. transformation and coarsening of the decomposition products of martensite.

The tempering process may be conducted at any temperature up to the lower critical temperature (Acl). Figure 29.17 illustrates the effect of carbon content of the martensite and tempering temperature on the hardness of carbon steels. 46 The specific tempering conditions that are selected are dependent on the desired strength and toughness. Nayar has recommended the following tempering conditions: 3~

�9 Heat the parts to 150-200~ to reduce internal stresses and increase toughness without a significant loss in hardness. This is often done with surface-hardened parts.

�9 To obtain the highest attainable yield strength and sufficient toughness, temper at 350-500~ �9 For optimum strength and toughness, heat to 500-700~ (Note: when tempering at high temper-

atures, 675-705~ precautions must be taken not to exceed the Ac~ temperature, above which undesirable austenite may be formed, which upon cooling would transform to pearlite. This is a particular concern for nickel-containing steels, since nickel depresses the Ac~ temperature. 47)

When steel is tempered in air, the heated oxide film on the surface of the steel exhibits a colour, known as a 'tempering colour' which is characteristic of the surface temperature. Table 29.14 provides a summary of characteristic surface temperatures for tempering and their colours. 32

In addition to the four steps shown, there is another step referred to as 'refrigeration' o r ' sub-zero' treatment. Sub-zero treatment is performed on steels to transform retained austenite to as-quenched martensite. Conversion of retained austenite in this way results in improved hardness, wear resistance and dimensional stability. Sub-zero treatment is performed using dry-ice or liquid nitrogen and involves cooling the steel to a temperature less than the Mf temperature of the steels which is typically between -30~ and 70~ (tool steels will elongate between about 0.5-2 ~m per mm of the original length during heat treatment 3 ). An immediate tempering step is required to remove residual stresses imparted to the steel by this process. Sub-zero treatments are not effective on steels that have been held at room temperature for several hours and therefore such treatments are typically performed immediately after hardening. 26 Although the general rule is to allow one hour per 25 mm (~ 1 inch) of the thickest cross-section, tool steels should be held at temperature for a minimum of two hours for each temper. 3

Heat t reatment o f steel 29-3 5

900 " I .... i ~ " i ] l I I t - ,,o i:,,,:'

As-quenched .. ~ 0- -"

~ o o - , ~ o o o J -

As-quenched #o ~ o " hardness / / ~

according to ___r / , . .

700 - Jaffe and Gordon t~" ,.o o o i o -

ooo / / / / o sov/O _ I , , , / / /Oo/~

/ / / / , / /

400 F / ~ / / O / 800~ O , ' s ,~o L,-- / /

E / O / ~ / O " 900~ O ~ O - 300 r , / ~ o ~ _ ~ . o -

i , " . - ~ ~ / o ' ~ 0 0 ~ o ~ - o - I-'. o ~ ~ ~ o ~%~00o~ o - - - - " ~ o.~ l - " / - ~ o ,,no~: o ~ - - - ~ o o ~ - - - " 200 F o , . . , / ' - ,..,~ " - ~ _ . ~ ~ - o - - , _ . . - - ~ , , _ . . . . v o . ' ~ O O L ~ . o " ' - - - I " . . o ~ / - - o ~ - I " 0 ~ "" ~ ~ 0 " " " -

[~. -- "8 ~ ~ 3 / - Tc = (TF - 32)/1.8

100 t i i ,I I I I I I I J 0 0 .2 0 .8 1 .0 0.4 0.6

Carbon, %

65

60

55

5O

45

40

35

30

25

2O

c) rr"

co (D r

" 0

"r"

Figure 29.17 Correlation ofcarbon content r martensite and hardness ot d(g]erent Fe-C alloys at d(/]brent tempering temperatures. (Tc and TF are the temperature in Centigrade and Fahrenheit re.v~ectivel.v. reproduced with permission of Metallurgical and Materials Transactions)

Table 29.14 TEMPER COLOURS OF STEEL

Abstracted from a much more detailed table in Thelning. 32

Temperature ( ~ C) Range o/tenq~er colourY

220-270 Straw yellow to brown 285-310 Purple to light blue 325-400 Various greys

Martensitic stainless steels and alloy steels that contain >_0.4% carbon and which exhibit M, temperatures of about 300~ are particularly susceptible to cracking, especially if they are through- hardened. 32 In such cases, cooling may be interrupted at about 80=C followed by an immediate temper at about 170~ to stop the formation of martensite. However, significant amounts of untempered


martensite remain after the steel is cooled to room temperature. Therefore, the steel must be tempered a second time, at the same temperature, to transform the hard and brittle, as-quenched martensite to a softer and more ductile, tempered martensite. This is one form of a 'double tempering' process.

Another form of double tempering occurs for high-alloy chromium steels and high-speed tool steels where significant amounts of retained austenite are transformed to martensite after tempering at about 500~ Such steels should be retempered to obtain a tougher martensitic structure. Generally, this second tempering step is performed at about 10-30~ below the original tempering temperature. 32

Steels that exhibit a high M~ temperature, _>400~ typically those that contain <0.3% carbon, form martensite which may be tempered during the remaining cooling (quenching) process. This is called 'self-tempering' or 'auto-tempering" and such steels are typically not crack-sensitive, particularly if the Mr temperature is > 100~C. 3-~

Typically, tempering times are a minimum of approximately one hour. Thelning has reported a 'rule of thumb' of 1-2 hours per 25 mm ( 1 inch) of section thickness, after the load has reached a preset temperature. 32 After heating, the steel is cooled to room temperature in still air. The recommended tempering conditions, in addition to the recommended heat treatment cycles, for a wide range of carbon and alloy steels are provided in SAE AMS 2759.

Tempering times and temperatures may also be calculated by various methods. One of the more common methods is to use the Larson-Miller equation discussed previously. The Larson-Miller equation, although originally developed for prediction of creep data, has been used successfully for predicting the tempering effect of medium/high alloy steels. -~ Bofors have reported that a Holloman- Jaffe constant of C = 20 was appropriate for all steels -~ but Grange and Baughman have reported that C = 18 should be used. 4s

Figure 29.18 49 shows that the Holloman-Jaffe constant varies with carbon content and desired 4s hardness. The incremental contribution to hardness of each alloying element in a steel may be

determined from Table 29.15. 49 The Vickers hardness (HV) is calculated by multiplying the concentration of each of the alloying elements (within the range shown) by the factor for that element at a given constant (C) and then all of these values are added together to provide the hardness (HV). The relationship between tempering time and the Holloman-Jaffe parameter (P) at different tempering temperatures is shown in Figure 29.19. 4~

The relationship between tempering temperature, time and steel chemistry has been reported by Spies .9.50

HB = 2.84Hh + 75 ( % C ) - 0.78(%Si) + 14.24(%Mn) + 14.77(%Cr) + 128,22(%Mo)

- 54.0 (%V) - 0.55Tt + 435.66

L 25

E

n 20

0.20

" ' , , ' ,

! , !

30 ~ ~ 1 i -

T---_ , , , r ' ~ - 'S" ' 350 OP~

, ~ ~ ~

~ ~ l 4 ,5"~'~o ~ ~ ~ ~ - - - - -

_

_! | o :t ~_ ~ I J , _ [L . . . . . . . [ - I ' i A i .L _ _ _

0.30 0.40 0.50 0.60 0.70 0.80

Carbon, %

Figure 29.18 Variation ofthe Holloman Parameter (Pt at df/tbrent hardness levels ( C .4S.~1 International, Materials Park, OH, used with pernfission )

H e a t t r ea tmen t o f s tee l 2 9 - 3 7

Table 29.15 FACTORS FOR PREDI(TING THE kI(KERS HARI)NFSS OF TEMPERED MARTENSITE

ASM International. Materials Park. OH. used with permission.

l-a('tors at indicated parameter (C) value ~

Element Range (%) 20 22 24 26 28 30

Mn 0.85-2.1 35 25 30 30 30 25 Si 0.3-2.2 65 60 30 30 30 30 Ni <4 5 3 6 8 8 6 Cr <1.2 50 55 55 55 55 55 Mo _<0.35 40 90 160 220 240 210

(20) 1 (45) I (8()) I (110) 1 (120) I (105) I V 3 <0.2 0 30 85 150 210 150

If 0.5-1.2~ Cr is present use this factor. 2 Note: the boron factor is 0. 3 May not apply if vanadium is the onh carbide former present.

1 0 0 - - - J ~I -. | [ r ! [ ! ! L - "~ - F. E " f - I l J i l ! i .

- : - [ I - [ J ! 1 l 1 i [ ! . 5o- �9 -1- I - l . . l I [ / ] ! I ! '1

- - - [ 1 [ [1 ! I I i I ! i 1 . . . . ojl

10- i ' [ ' 1 ] [ i , . ! ! I !~J~~ I - �9 , :

I - l - i { t i , * [ I 5~_ ] I - J - I [ ! ! ! [ 1 ] iT/ I [ I ]

] ] ] - J - [ /_ 1! / t i I i ] / [ I / - 1 -![ I] ![ I I~ I [ l i i

o . . . . r 't I ! ' ] [ i - " I ! 1 [ 1 1 ~ , ~ 2] - ! - - i i11 ! ! J ~ ~I[ 'J[ I

o -1 ~ -1 - ~ ! i l ]1 I ' l ] ! / oo 1 ] . . . . , i ~ { 6 0 I -' : I ~ I : t { I { " ; { I

.E. - : , "~ ] ~ ~ 1 [1 [ 1 ] , [ ~ ; I ' . I ' " i - 4 0 ~ ; 0 . 5 - ]I ~ ] _ I [i I ! J [ ~ ! ! i/ ]

_ _ _ - i . ~. 1 ( 1 )

~- ] [ ! ! ~ I ! !1t i ] 1 ' i - 2 o - ~ ' : 0

o.~] ; I I-~ ; i~ /! l i l ! l ] i [ ; i i - l l 1 ] /_ I1! ~ ] : 1 1 1 ! i l : i ; - ~ o

0 . 1 ] l - ! I l ],':- I - I ~ ~ !I" | ].~ : ~ ! ~ ' -

, ~ i l i if! t ; l l - ~ ' ' I~ �9 l ~[ I ! i l l I 1 ]

0 . 0 5 ~ l ] ] ]!! ]! [ ; ! i f : ] ] ~ [ ] ] 3 E

_ -ii 1 i !] I 1 111 ; 1 ! : ; i -

15 20 25 30 3 5 Parameter

(~ + 460) (18 + Log of time in hours) x 10- 3

Figure 29.19 Time-temperature ~'er~ux t t ( ) / I . m a n Parameter chart tin ( = 1,'~ ( c . - I S . i t International. MateriaL~" Park. OH. used with permi.~.~ion )


250

200

�9 ~ 150

( / )

,,..,

09 100

50

Tempering temperature, ~

210 390 570 750 930 1110 1 1 I 1 I [ ]

/

Tensile strength

.,,..

- R A

E l o n g a t i o n _ . ~ m,

I I.,. I I l t _

1 0 0 200 300 400 500 600

100 d 500

O

80 = m .o "1-- 400

60 ~ L r " O

,'- m 300

40 c 0

c 200 O 20

Tempering temperature, ~ (a) (b)


210 390 570 750 930 1110 . . . . . . . .

1 1 1 l l 1

_ 1 0 0

Brinell

- 80

60

40

- 2 0

I l I l I_ l 0 100 200 300 400 500 600


,,.2. t - O

0 t -

v

+6

E

t-~ L r t -

O

Figure 29.20 Loss in room-temperature toughness due to temper embrittlement of oil-quenched wrought Ni-Cr-Mo (AIS14130) steel. section si-e = 13 mm (0.5 in. )

where HB is the Brinell hardness after hardening and tempering, Hh represents the Rockwell C (HRc) hardness after hardening and Tt is the tempering temperature in ~ This equation was developed for the following conditions: Hh = 20-65 HRc, C = 0.204).54%, Si = 0.17-1.40%, Mn = 0.50-1.90%, Cr = 0.03-1.20% and Tt = 500-650~ (932-1 202~

The German DIN 17021 standard provides a relationship between the as-quenched hardness (Hh) and the as-tempered hardness (H,) :~

Hh(HRc) = ( T t / 1 6 7 - 1 .2)Ht- 17

where Tt = tempering temperature in :C. Temper embrittlement may occur for some steels if they are tempered below 595~ which is

observable by a reduction in notch toughness as illustrated in Figure 29.20 for AISI 4130 (Ni--Cr- Mo), steel which was tempered at 260-370~ ~7 However, temper embrittlement may be avoided by tempering at higher temperatures, with subsequent quenching, to minimise the time the steel will

17.47 spend in the intermediate temperature range. Tempering may be performed in convection furnaces, salt baths or even by immersion in molten

metal. Induction tempering and flame heating are also used but will not be discussed here. Table 29.16 provides a comparative summary of the different heating media. 5~ Of the different tempering systems shown, convection furnaces are the most common and it is important that they be equipped with fans and/or blowers to provide uniform heat transfer, when heating the load. Typically, convection tempering furnaces are designed for use in the 150-750~ range.

Salt baths may also be used for various heating processes over the temperature range 150- 1 320~ 19 and they provide relatively rapid heat transfer, compared with convection furnaces, although the actual use temperature is dependent on the composition of the salt bath. A comparison of heating rates between salt baths and muffle furnaces is provided in Table 29.17. 32

19 Sinha has classified salt baths into three groups:

�9 Low-temperature salt baths may be used from 150-620~ These baths are of two types: a binary mixture of equal parts of potassium nitrate and sodium nitrite, which may be used for heating to 150-500~ and binary mixtures of potassium nitrate and sodium nitrate which may be used for heating in the range of 260-620~ In addition to tempering, these baths may also be used for cooling. It is essential however, that the baths are not contaminated with cyanides, organic compounds or water!

H e a t t rea tment o f s tee l 29-39

Table 29.16 TEMPERING TEMPERATURE RANGES A('HIE\~BLE WITH DIFFERENT TEMPERING EQUIPMENT

ASM International, Materials Park, OH. used vdth permission.

Temperature range Equipment O,pe ~ C ~ F ~.~e conditions

Convection furnace 50-750 120-1 380

Salt bath 160-750 320-1 380

Fluidised beds 100-750 212-1 380

Oil bath <250 <480

Molten metal bath > 390 > 735

For large volumes of similar parts: lbr variable loads. Temperature control more difficult.

Rapid uniform heating. Low to medium volume. Should not be used for complex, hard to clean parts.

Broad range of heat transfer rates are possible by va~'ing the choice of fluidising gas. gas velocity, bed temperature and the bed particle size. More energy efficient than convection furnaces and they provide a safe and ecologically friendly alternative to salts and lead. with similar heat transfer rates.

Good if long exposure times are desired: special ventilation and fire control are required.

Ver~' rapid heating: special fixturing required: molten metals may be toxic (Pb baths).

Table 29.17 HEAT-UPTIMES FOR IOOmm DIAMETER BARS TO 950-1000~C

Abstracted from a more detailed table in Thelning. 3:

Salt bath .lh(/]h'.liovmce (min) Iminl

8 60

�9 Medium-temperature neutral baths which are suitable for use over the range of 650~ O00~ These baths are binary or ternary mixtures of the following salts: potassium chloride, sodium chloride, barium chloride or calcium chloride. Two examples of typical binary compositions and working temperatures include: NaC1 (45%)/KC1 (55%) which is suitable for use at 675- 900~ and NaC1 (20%)/BaC12 (80%) which is suitable for use at 675-1 060~ BaCle, if used at 100%, has a relatively narrow use range of 1 025-1 325:C. -~2 An advantage of these baths is that when they are freshly prepared, the steel surface will be clean without surface carburisation or decarburisation.

�9 High-temperature salt baths are used in the range of I 000-1 300:C and they typically contain mixtures of barium chloride, sodium tetraborate (borax), sodium fluoride and silicates. These baths may decarburise steels, as oxides build-up after use.

Molten metal, most typically lead, has been used in the past, but due to its toxicity, its use is now restricted to various heating operations where its outstanding heat transfer properties are essential. Lead baths are used from 327~ which is the melting point of lead, up to 900~ 1'~

Fluidised beds are formed by passing a gas through solid particles such as aluminium oxide and silica sand, which causes the particles to behave like a bubbling liquid. The particles are generally inert and do not react with metal parts but act to facilitate heat transfer between the fluidising gas and the part being processed.

A broad range of heat transfer rates are possible over operating temperatures which may range from 100-1 050~ (212-1 920~ with fluidised bed furnaces for tempering operations typically ranging from 100-750~ (212-1 380~ (see Table 29.16). Fluidised bed furnaces are not only more energy-efficient than convection furnaces but also exhibit heat transfer efficiency similar to salt baths and lead pots, without the health and environmental safety hazards commonly associated with these systems.

29--40 Heat treatment

60

56

52

48

44

40

,--- 36 ._ E �9 " 32 E I- 28

24

20

16

12

Steel diameter, in.

0 0.98 1.97 2.95 3.94 w 1 l i ~ . . . . . !

How shapes affect heating speeds Radiant __.] -- Shape Speed tactor atmosphere | _ D.x.,,.-~---- ~ Long cylinder 1 fumace/__J

(din. D) / q

__ Tw,,~=~z:=~ Long square 1 i / _...~ D ~ ~ (D• D) Vacuum/ [

-- ~ , , ~ : ~ ~ Long rectangle 0.7 .~ / --.4 .'" . - /

i !-- - ~ Long rectangle 06 "" " / --" L D I H 3 D _ ~ (D • 3D) " . " ' V ~- / , Infinite plane 05 ' / / - / ' L

D' (very wide, " ." / ~ t thickness D) o'" D ~ Sphere 1.5 , " / / ' "

I" ~ (d in , D) .,'/" / - -

/ I ' f " - ' - ~ Cube ' = "" / / Radiant D~i I I D • D >~ D . . . . " / / ' fluidised - - [ ~ u c ~ / / . / ~;;~ . . . .

-- , ' / / " (Firewall) --

-- ~/ ,o*" ./i,.s,, ,, ~ ,. (Fireb~ id ~ "*s~d'* *" / .- " " -" Salt

_ . / -

: ...... .......

o 1 1 1 1 1 0 25 50 75 1 oo

Steel diameter, mm

Figure 29.21 Recommended heating times/or various heating media and.shapes

Figure 29.21 provides a comparison of the relative heating rates that can be achieved with these �9 ~7 different heating media.-

29.2.9 Austempering

Austempering requires that the cooling process be fast enough to avoid the formation of pearlite, as illustrated by the cooling curve for an austempering process, which is superimposed on to an ITD (TTT diagram) of a steel in Figure 29.22Afl The steel is cooled below the nose of the pearlite transformation curve for the steel being quenched and just above the M~ temperature, in a molten salt bath. The steel is held at this temperature until the transformation from austenite to bainite is complete. Upper or lower bainite may be formed depending on the molten salt temperature. Austempering eliminates the volumetric expansion due to martensite, which helps to eliminate cracking and provides greater toughness and ductility at a particular hardness�9

In practice, austempering is performed by heating the steel to 790-870~ and then quenching into a molten salt bath at a temperature just above the M~ temperature (260-400~ The steel is kept at the molten salt bath until the transformation is complete, at which time it is cooled to room temperature in air. 3~

The suitability of steels for austempering depends on: sz

�9 The location of the nose of the pearlite start (P~) curve on the ITD and the ability to cool the steel sufficiently fast to avoid pearlite formation�9

�9 The time required for complete transformation from austenite to bainite, at the austempering temperature.

�9 The Ms temperature of the steel.

L

?

A3

Ms |

A1


(a)

TT

(b)

Process steps 1. Austenetise 2. Quench 3. Hold just above M s 4. Bainitic product 5. Hold below M s 6. Martensitic product

A

|

@

=, ._

v

log t

A3

A1

v

log t

Figure 29.22 Illustration ~?f (a) austempering. (h) martempering (redrawn. with mod(fications, after relerence 9. p. 456--courte~:v ~?/.~huz'el Dekkep: Inc. )

29.2.10 Martempering

Martempering (also known as marquenchingl is a cooling process used to minimise distortion and cracking and to reduce the amount of residual stress formation of a part upon cooling. Typically, martempering involves austenitising steel at a temperature above Ac~ (815-870=C) and then quenching into hot oil (marquenching oil) or molten salt at a temperature just above the M~ temperature (usually at a temperature ranging from 260-400=C) followed by holding at this temperature until the entire steel workpiece is at the same temperature, at which time the workpiece is cooled to room temperature, at a rate sufficient to achieve a fully martensitic structure. Note that if the part is held at the martempering temperature too long, undesirable formation ofbainite may occur. -~ Usually the quenching step is performed in air which minimises thermal gradients and related residual stresses throughout the part while cooling. This process is illustrated in Figure 29.22B. ~ After quenching, the part is tempered per specification.

A variation of this process is to cool the part to a temperature intermediate between the M~ and MF temperature, until the temperature at the surface and core have completely equilibrated, at which


Table 29.18 TEMPERATURE RANGE OF MARTENSITE FORMATION OF SEVERAL CARBON AND LOW-ALLOY STEELS

Steel grade AISI

50% .Plartensite M f--99% Martensite Ms temperature /brined formed

~ ~ ~ ~ ~ ~

1030 343 650 293 560 232 450 1065 277 530 219 425 149 300 1090 219 425 157 315 82 180 1335 343 650 293 560 232 450 2340 304 580 293 560 207 405 3140 335 635 288 550 221 430 4130 380 715 343 650 288 550 4140 343 650 299 570 227 440 4340 288 550 249 480 188 370 4640 343 650 299 570 249 480 5140 343 650 299 570 232 450 6140 327 620 293 560 232 450 8630 366 690 332 630 277 530 9440 330 625 282 540 207 405

Table 29.19 TYPICAL USE TEMPERATURES FOR MARTEMPERING OILS

L'se temperatures

Viscosi~ Minimum flash point Open air Protective atmosphere at 40~ SUS ~ :F ~C :F ~ ~

250-550 220 430 95-150 200-300 95-175 200-350 700-1 500 250 480 120-175 250-350 120-205 250-400

2 000-2 800 290 550 150-205 300-400 150-230 300--450

time the steel is removed from the bath and air cooled or placed directly in the tempering furnace. Table 29.18 provides a summary of Ms and Mr temperatures for a variety of common steels. 17

Either martempering (hot quenching) oils or molten salt baths may be used for martempering. Molten salts were reviewed briefly above. Martempering oils will be discussed briefly here. Martempering oils are used at temperatures up to 205~ sometimes as high as 235~ Molten salts are used for martempering operations performed at 204-400~ 52 The oils are usually formulated with solvent refined petroleum oils with a very high paraffinic fraction, to maximise oxidative and thermal stability, which is further enhanced by the addition of antioxidants. Accelerated and non- accelerated martempering oils are available. Typical temperature ranges for commercially available martempering oils are summarised in Table 29.19. 53 Because martempering oils are used at relatively high temperatures, a protective, non-oxidising atmosphere is often employed, which allows the use temperatures much closer to the flash point, than is generally recommended for open air conditions.

More hardenable alloy steels are martempered more often than less hardenable plain carbon steels. Steels that are most often martempered to full hardness include: AISI 1090, 4130, 4140, 4150, 4340, 4640, 5140, 6150, 8630, 8640, 8740, 8745. Carburised steels from 3312, 4620, 5120, 8620 and 9310 may also be martempered. Plain carbon steels 1008 through 1040 are insufficiently hardenable to be martempered. Thin sections of some steels of borderline hardenability such as AISI 1541 can be martempered. 52

29.2.11 Carburising

THERMODYNAMICS OF THE CARBURISING PROCESSES

Carburising systems involve reversible chemical reactions of the form:

x A + y B ~ q C + r D

for which the equilibrium constant (I~q) is:


Keq (pq r x ,' = PD)/(PA Pi3)

where PA, PB, PC, PD are the partial pressures of reactants A and B and products C and D and x, y, c and d represent the stoichiometry of the reaction. In the event that product C is dissolved in the steel, the activity of C in the steel (ac) would be substituted for Pc.

These reactions attempt to approach equilibrium involving the atmospheric constituents and the substrate material (steel). At equilibrium, the change in Gibbs free energy (AG) is zero by definition. Hence:

AG = A G ~ + RT In Keq = 0

Thus, at equilibrium, the equilibrium constant, K~q, is given by:

A G ~ = - R T In K~q

where:

AG~ free energy of the reaction [kJ mol-I], R--Gas constant (8.314 J K- 1 mol- 1 ), TmAbsolute temperature [K].

A carbon-rich atmosphere that is capable of carburising steel objects, when held at an appropriate process temperature, is known as a controlled carburising atmosphere. Decarburising/carburising is governed by the following reversible reaction:

C(iny) + CO2(g) ~ 2CO(g)

which has an equilibrium constant of KB = p~-o/(Pco, ac ) When using methane as the process enrichment gas, the reaction involved is:

CH4(g) ~ C(iny) + 2H2(g)

for which the equilibrium constant is K xl = (p~l. ac)/pcii~ _

As well as the following:

C(in y) + H20(g) ~ CO(g) + H2(g)

with an equilibrium constant K~- = Pco Pll-/(P~t:o a~.), where K~, KM, Kw are the equilibrium constants of the reaction concerned, ac is the carbon activity, Pco, Pco., Pit_., PH.O, PCH~ are the partial pressures of the species shown.

Thus, by selection of an appropriate atmosphere, carbon may be introduced into the austenite as a solid-solution. Alternatively, free carbon may be present in the furnace atmosphere, in the form of soot (graphite). If the atmospheric carbon concentration is high enough, graphitic soot will precipitate out of the atmosphere and begin to concentrate on the work piece, the walls of the furnace and mechanical parts within the furnace, such as the rails, hearth, rollers, burner tubes and elements.

The equilibrium constants for the reactions shown above are:

lOg KB -- --8 750/T + 9.022

log KM ---- --4 768/T - 5.767

log Kw --- - 6 908/T + 7.457

KINETICS OF THE CARBURISATION PROCESS

Carbon deposition at the steel surface

Consider now the rate at which carburisation occurs. At the steel/gas phase interface the carburisation reaction will be dependent on the difference between the carbon potential (carbon activity) in the atmosphere and at the steel surface. Carbon will diffuse from the atmosphere to the steel surface


when the activity of the carbon in the atmosphere is higher than the activity of carbon in the austenite surface, which is dependent on furnace temperatures and the initial carbon concentration in the steel. Thus, the number of carbon atoms, N, that penetrate a surface area S, in time dt, is dependent on the difference between the carbon activity in the gaseous medium c~cg and the carbon activity on the steel surface C~cp:

N - /~(e,c~ - ~p)

S dt

If the carbon concentration is substituted for activity then:

M - /3(CA - C s )

F dt

where:

M--mass of carbon deposited (g), F--surface area (m 2) penetrated by carbon of mass M,

dt--time of carbon penetration (s), CA--carbon potential of the atmosphere (g m -3 ), Cs--surface carbon content (g m -3 ),

fl--surface reaction rate constant (m s- ~ ).

The carbon transfer coefficient is defined in European norm EN 10052:199354 as: 'the transfer of a carbon mass from the carburising medium to the steel surface, through individual carbon molecules to the steel surface by the difference of potentials between carbon potential of medium, and the steel carbon potential at the surface at any given moment. The mass transfer stream of carbon from atmosphere to steel is quantifiable in the following formula':

Carbon flux (in gs -I ) = flc(Cp - Cs)F

where:

tic--carbon transfer coefficient (m s-l), Cp---carbon potential of atmosphere (g m- 3 ) Cs---carbon content of steel surface (gm -3)

F--surface area (m 2)

which is identical to the previous equation, apart from the change in nomenclature. Assuming that steady-state conditions are established, the rate of carbon transport to the steel

surface can be described by means of Fick's first law of diffusion, which states, using nomenclature appropriate to the diffusion of carbon, that:

OC Jc = -Dc-o--~x

In this equation, Jc is the carbon flux, i.e. the amount of carbon passing through 1 m 2 of reference plane per second (gm -2 s -1 ). C represents the carbon concentration (gm -~) and x is distance (m), so that the one dimensional carbon concentration gradient is (aC/ax). Note: any convenient measure of the amount of carbon can be used in place of grams in this equation, but a consistent unit must be employed for the amount of carbon in the C and J terms. Dc represents the carbon diffusion coefficient (diffusivity) in the medium in which the carbon is diffusing (and has units of m: s -1 ). The value of Dc will depend strongly on the carburisation temperature.

Re-expressing the carbon flux in Fick's first law in units of g s-1 yields:

dC Carbon flux (in g s- 1 ) _ Dc F

dx

Given continuity of carbon mass flow from atmosphere to steel and then into the steel surface, it is then possible to equate the two expressions for the carbon flux, which yields the following mass transfer formula:

Dc Cp - Cs

/ 3 c dC/dx


The value of the carbon transfer coefficient tic (and hence the capability to transport carbon to the steel surface) depends on the source and hence the carbon potential of the medium used for carburising. From the literature, the smallest carbon transfer coefficient tic is 5.4 • 10-~o m s -~ , for the mixture CO-CO2. A somewhat larger value of 1.79 x 10 -~ m s-1 has been reported for a mixture of CH4-H2-H20. Mixtures of CO-CO_~-CH4-Hz-H~O exhibit higher tic values: 3.1 x 10-" ms -l for a mixture of 50% CO-CO_~ and 50% CHq-Hz-HzO. An industrial endothermic atmosphere composed of fuel gases exhibits a tic value of 1.2 x 10-- m s -~ .

Diffusion of carbon into the bulk

In the last section, Fick's first law was used to describe the steady-state diffusion of carbon to the surface of the steel, such that a fixed carbon concentration Cs is established at the surface of the steel. Within the steel, however, non-steady state carbon diffusion occurs as the case develops and so it is necessary to employ Fick's second law of diffusion, i.e.:

OC O2C 3--T = D~. 0x 2

In this equation, t is time (in seconds) and D~'. is the diffusion coefficient of carbon in austenite at the carburisation temperature. Other terms are as defined previously (although, in the present case they refer to the carbon concentration within the steel}. This equation does not permit a general analytical solution. However, the following result can be obtained for the carbon concentration as a function of distance and time, C(x, t}. during carburisation of a steel with an initial carbon concentration CI prior to carburisation:

In this case, x = 0 is defined as the surface of the steel in contact with the carburising atmosphere, which has a fixed carbon concentration C s. This equation may be solved with appropriate error function (erf) tables, or a spreadsheet that supports the error function. Alternatively, if all that is required is an order of magnitude estimate of the case depth, then the following approximation can be employed:

Case depth ~ ~ c ' - t

In both of these equations, the assumption is that the rate controlling step governing the formation of the case is the solid-state diffusion of carbon in austenite. The value of D~i for use in either of these equations can, in turn, be calculated using the Arrhenius equation:

" [ D~ = (D0)~ exp - ~ - ~

Notice that both the pre-exponential term (Do)~'-, sometimes called the frequency factor (with units of m 2 s -I ) and the activation energy for diffusion, Q~i (with units ofkJ mo1-1 ), are material specific to the diffusing solute (carbon} and matrix (austenite) involved. In this equation, R is the gas constant (8.314 J K-1 mol- ~ ) and T is the absolute temperature (in K) at which carburisation is performed. Table 29.20 contains tabulated values for D~i as a function of temperature. See Chapter 13 for diffusion data.

CLASSIFICATION OF CARBURIS1NG PROCESSES

The most common processes that are encountered industrially are:

a. Pack carburising--The pack carburismg process is typically conducted by surrounding the steel in a pit furnace or steel box furnace with granules of charcoal or charcoal plus coke. An 'activator' for the charcoal, such as barium borate (BaBO3) is added to facilitate the release of CO2, which then reacts with excess CO~ to form CO which in turn reacts with the low- _ carbon steel surface to form carbon, which diffuses into the steel. Pack carburising is conducted typically at 920-940~ for 2-36 hours.


Table 29.20 EFFECT OF TEMPERATURE ON THE DIFFUSION COEFFICIENT OF CARBON IN AUSTENITE IDa')

Temperature

~ ~ D~'. (m 2 S - I )

760 1 400 7.8 x 10-13 788 1 450 1.2 • 10 -12 816 1 500 1.9 x 10 -12 843 1 550 2.8 x 10 -12 871 1 600 4.2 • 10-12 899 1 650 6.0 • 10 -12 927 1 700 8.6 • 10 -12 954 1 750 1.2 • 10-1 982 1 800 1.6 x 10-1

1 010 1 850 2.2 x 10 -I 1 038 1 900 3.0 x 10-1 1 066 1 950 4.0 x 10-1 1 093 2 000 5.2 x l O-I 1 121 2050 6.8 x 10 -1

b. Liquid carburising--Liquid carburising is conducted typically in internally or externally heated molten salt pots containing a cyanide salt such as sodium cyanide (NaCN). There are generally two types of liquid carburising processes. One type is a low temperature process (840-900~ which is conducted when low case depths of 75 to 760 ~tm (0.003-0.03 inches) are required. The second liquid carburising process is conducted at a high temperature (900-950~ when case depths of 760 to 3 050 ~tm (0.03-0.12 inches) are desired. In either case, process times may be 1-4 hours.

c. Gas carburising--Currently, the most common carburising process is gas carburising which may potentially be performed with any carbonaceous gas such as: methane, ethane, propane, or natural gas. Carburising times o f4 -10 hours are typical. The carburising temperature is greater than the upper critical temperature (in the austenite transformation region, >954~ Case depths are typically less than 1.3 mm (0.05 inches). The conventional gas carburising process allows measurement of the gaseous carbon activity within the furnace process chamber, using the shim test method, dew point test method, CO/CO2 test method and oxygen probe.

d. Vacuum carburising--Vacuum carburising is a clean method used to introduce carbon into the surface of the steel and also prevents grain boundary oxidation. Vacuum, or low-pressure, carburising is carried out in a vacuum furnace at pressures below normal atmospheric pressure. The principle of carburising is exactly the same as that of the gas carburising process, the main difference being the use of subatmospheric pressure.

e. Ion carburising--Carbonaceous gases such as methane have been used for vacuum carburising because of their widespread availability and current use in gas carburising. Although methane is reactive and controllable when used with endothermic atmospheres, methane alone is extremely stable, even at elevated temperatures. Therefore, when methane is used in vacuum carburising, relatively high pressures of 33 to 53 kPa (250-400 torr) are required. Furthermore, carburising processes conducted at these relatively high pressures produce significant amounts of carbon sooting. These deficiencies however, may be overcome by using an ion process. Methane ionisation is produced with a high-voltage (approximately 1000 volts) at a relatively low pressure of 1.3 kPa ( 10 torr). When methane ionises by this 'ion carburising' process, a reactive gas blanket is formed in close proximity to the workpiece, without concurrent soot formation. Other hydrocarbons can be used similarly.

f. Fluidised bed carburising--Steel carburising processes may also be conducted in fluidised bed furnaces. Various atmospheres may be used, including conventional endo gas/hydrocarbon mixtures or nitrogen/methanol/hydrocarbon mixtures. Depending on the carburising atmosphere, fluidised bed temperatures of 850-975~ for 30min to 3 hours may be used. Case depths of up to 700 tzm are typical.

Table 29.21 provides a comparative summary of the relative advantages and disadvantages of these processes.

Table 29.21 COMPARISON OF CARBURISING METHODS

Heat treatment o f steel 2 9 - 4 7

Description of method Disadvantages Advantages

A. Carburising with solids (pack carburising) 1. The basic component of the

carburising medium is ground wood charcoal (around 3-5 mm granules), which is mixed with carbonates of barium, sodium, calcium, lithium or potassium.

2. The temperature used for carburisation is about 900~

3. It is necessary to place the components into a steel box with a spacing of 25 mm between the components. The box design is very simple, although this does require a lid which can be sealed with clay to contain the liberated gas.

4. Starting the reactions during heating can be accomplished by: �9 Burning of a small quantity

of the wood charcoal by the introduction of oxygen.

�9 Reaction of an activator with wood charcoal.

5. After the carburising, the carburising boxes can be cooled down within the furnace. Alternatively, the boxes can be removed from the furnace whilst hot and air cooled.

A long heat up time is necessary to reach the process temperature and to achieve temperature uniformity, throughout the box.

Decarburisation of the steel surface will occur if the components are allowed to air cool without protection, or removed from the process box.

It is difficult (but possible) to harden directly from the carburising box. It is usual to allow the box to cool down and then reheat to the required austenitising temperature.

The method is not reliable in terms of repeatability and cannot be controlled accurately. This is a slow production method.

Grinding is necessary, after the procedure, due to a slight potential for surface porosity.

B. Liquid carburising 1. A mixture of molten salts is 1.

the carburising medium. Usual mixtures are carbonates or chlorides and cyanides of alkaline metals, sometimes with an addition of SiC. A typical mixture would be: 75% NazCO3, 2. 15% NaCI, 10% SiC.

2. The temperature used for carburisation is usually between 900~ and 950~

3. The steel components are 3. placed directly into the molten salts after pre-heating. A cocoon of salt will immediately 4. adhere to the steel surface, thus offering some thermal protection. 5.

4. In the presence of iron, the cyanide salt decomposes at the high temperature of the bath and cyanate (CN) is liberated. which further decomposes to provide carbon to diffuse into the steel.

There is a large amount of sludge that collects in the bottom of the salt bath. It is mandatory to clean out the sludge on a frequent basis.

High operating costs, due to post cleaning and effluent disposal. The process is labour intensive and involves long pre-wiring times for the components.

Carburising stop off is difficult.

Components need a long pre-heat time.

The probabili~" of distortion remains.

1. Low capital equipment cost. cost.

2. Simple procedure.

3. Low operating costs.

1. Possibility of direct hardening.

2. Uniform carburising of clean surfaces.

Elimination of steel boxes that are used in pack carburising procedures.

It has been possible to accelerate the process considerably by applying an electrolytic method.

(continued)

29--48 Heat treatment

Table 29.21 COMPARISON OF CARBURISING METHODS--continued

Description q['method Disadvantages Advantages

C. Gas carburising

1. Atmospheres for carburising 1. are produced in special generators which produce a process gas from natural gas blended with air. The generators are known as endothermic generators. The natural gas will contain a 2. large proportion of methane, plus lower concentrations of other hydrocarbon gases. The air that is mixed with the natural gas will contain moisture which will assist in controlling the carbon potential of the endothermic gas.

2. The temperature used for 3. carburisation is usually between 870-950~

3. The furnace heating system will assist the gaseous atmosphere to ensure good temperature uniformity within the furnace process chamber. In addition to this, the carbon potential is usually very uniform throughout the process chamber.

D. Vacuum carburising

1. The process is conducted at sub-atmospheric (partial vacuum) pressures. Process gases of methane, propane, or acetylene are introduced into the process chamber.

2. Atomic carbon is generated as a result of the break-up of the process gases. The process of vacuum the carburising involves following stages: �9 Saturation of the

atmosphere with carbon at the process working pressure.

�9 Diffusive transportation of excess carbon into the steel surface in high vacuum.

3. Single chamber vacuum furnaces can be used in anv configuration. Also, front or rear cooling chambers can be fitted, that can facilitate controlled cooling of the processed batch after austenitising. High pressure gas quenching can be accomplished when using blended gaseous mixtures of either nitrogen and helium or nitrogen and hydrogen. Blended gas quenching can (depending on the gas blend and delivery pressure), equal the quench speed of oil.

Limited speed of diffusive satiating, resulting from limitations of the furnace construction as well as the carbon potential of previous carburising atmospheres.

Carburising atmospheres will contain oxygen in the form of moisture, which will cause intergranular oxidation and also create the potential for grain boundary corrosion. This will cause a deterioration of the fatigue strength of the carburised case.

Considerable emission to atmosphere of harmful substances (oxides of carbon and heated quench oil effluent resulting from quenchingt. It is advisable to install fume extraction systems to ensure adequate shop ventilation.

Oil quenching may lead to distortion and possibly the risk of cracking.

1. High capital cost of equipment.

2. High operating costs.

3. Difficult process control in temas of determining and controlling the atmospheric carbon potential.

Easy to change the carbon potential of the furnace atmosphere, simply by adjusting the enrichment gas in relation to either the moisture present in the atmosphere or by the presence of free oxygen.

Small waste of energy and economy of time.

Parts are relatively clean except from oil quenching if the furnace has an oil quench system, such as is seen on an integral quench furnace.

Possibility of hardening directly after carburising. However great care must be taken when considering this.

Fast process times at conventional carburising temperatures.

Advantage can be taken of high temperature carburising, up to 1 075~ Since carburising process temperatures approximately 150~ higher than for conventional gas carburising temperatures can be used. this will result in a shorter gas phase transportation time, resulting in much faster carburising times. The superficial concentration of carbon as a result of an unequal process of break-up of hydrocarbons is, as a rule, very high. Both of these factors will accelerate considerably the diffusive saturation.

The carburised layer shows the best mechanical properties.

Clean finished work surfaces that do not require post cleaning.

Environmentally friendly vdth no toxic gas emissions. Also, lower volumes of effluent gas than for conventional processes.

(continued)

Table 29.21 COMPARISON OF CARBURISIN(i METH()DS--('otltimted

Heat treatment o f steel 2 9 - 4 9

Description of method Disad~'an rages .4dvantages

D. Vacuum carburising--contbzued

E. Ion carburising

1. The process depends on placing the steel in a vacuum chamber in a low pressure hydrocarbon atmosphere, with the simultaneous application of a high voltage, on heating.

2. The furnace wall is the anode whilst the work load sits on the furnace hearth at the cathode's potential. A voltage is applied in the region of 450 to 800 volts (dependent on the chamber pressure and the workload surface area), which will cause a glow discharge in the process chamber. With the work load at cathode's potential, the process gas is immediately ionised and carburisation will begin to take place.

3. The steel surface does not act as a catalyst as occurs with the more conventional carburising procedures. However, the process gases used are still methane, propane or acetylene, along with nitrogen and hydrogen.

1. High capital investment.

2. Analysing the carbon potential within the process chamber is difficult. Thus, determination of the carbon potential must usually be accomplished prior to each batch, using data acquired from previously carburised loads.

6. Mechanical handling equipment can easily be installed into the system for part transportation within and outside of the furnace.

Effective. energy-saving process.

7. Hydrocarbons are the .

exclusive carrier of carbon. Therefore there is no risk of grain boundary oxidation.

8. Cooling of the work load can be accomplished, if necessary, under nitrogen. This will eliminate the need for post washing.

9. The quench gas direction can be manipulated to suit the part geometry, thus reducing the risk of distortion. In addition, the gas floyd rate can be adjusted v~hen using a two speed gas circulation drive motor. Thus, the risk of distortion is less than when quenching into oil.

It is possible to control the thicknesses and structure of diffusive layers.

It has been possible to carburise and successfully treat (on a repeatable, continuous basis) stainless steels. heat-proof and acid- resistant steels.

Case uniformity is excellent, irrespective of the part geometry.

(continued)

29 -50 Heat treatment

Table 29.21 COMPARISON OF CARBURISING METHODS--continued

Description of method Disadvantages Advantages

E Carburising influidised beds

1. A fluidised bed is created usually as a result of the activation by a gas passing through a bed of particles such as sand or aluminium oxide. The particles of the bed are kept in suspension by hot process gas passing upwards and through the particles.

2. Heat transfer is rapid in a fluidised bed.

The fluidised bed furnace can be heated directly or indirectly. Heating can be electrical or by gas, with the enrichment gases being added with the heating gas.

The method of operation is exactly the same as is with the salt bath method of heat treatment, with the parts simply being immersed into the fluidised particles.

1. No major problems. 1. Simple, but efficient furnace design.

2. Low operating cost.

3. Ease of operation.

4. Not labour intensive.

5. The parts are not wet as in a salt bath. There is no slag to handle, or desludging of the bed to perform.

6. Good heat transfer up to the process temperature.

7. It has been possible to apply direct hardening of components by reducing the carburising temperature to the appropriate austenitising temperature.

STEEL GRADES USED FOR CARBURISING

Construction alloy steels that can carburised successfully are many and varied. When selecting a steel for carburising, it is important to consider machinability, resistance to overheating, susceptibility to deformation during thermal processing, hardenability relative to the cross-sectional size and geometrical features, in addition to mechanical strength, not only in the case but also in the core.

Typically, steels selected for carburising contain <0.25% carbon. The alloy composition is selected to provide case and core hardenability. Plain carbon steels may be carburised but the carburising response is limited owing to the lack of alloying elements. This is illustrated by the selected listing of carburising steels provided in Table 29.22. 55

Case depths of 75 to 6 350 ~tm (0.003-0.250 inches) with surface hardnesses o fRc = 58-62 are usually specified. An 'effective case depth' is typically defined, which may be determined in various ways. One method is to measure the case depth on a metallographic sample of the part, or of a test bar, by determining the microhardness at various depths from the surface. The required desired case depth is govemed by the end application, such as those shown in Table 29.23. 55

Table 29.24 provides a summary of the common carburising grades of steel used internationally.

HEAT TREATMENT AFTER CARBURISING

After carburising, austenitisation is necessary. If the workpiece is quenched directly from the carburising temperature, retained austenite will form, along with a coarse grain size (depending on the carburising time at temperature). Post-carburisation heat treatment is necessary to provide high surface hardness, that will resist abrasion and wear. Depending on the requirements of the carburised part, the post heat treatment temperatures are selected to provide not only a high surface hardness but also the appropriate core strength. Some options for heat treatment after carburising are shown in Figure 29.23.

The hardening methods summarised in Figure 29.23 are the most commonly used. However, the following procedures are also used, although less commonly:

,, Hardening after pearlitic transformation will provide greater energy efficiency and will assure grain size reduction.


Table 29.22 SELECTED LIST OF STEEL COMMONLY USED FOR CARBURISING AND THEIR FEATURES

29-51

Steel grade Features and bene/~ts

4620

8620

4320

4820

9310

Relatively low cost, chromium nickel molybdenum steel. Used only where nominal hardenabilitv and core response are required.

Most commonly specified steel for carburising. Excellent carburising response, with good hardenability for most section sizes.

Higher hardenabilitv for improved core response in thicker cross-sections.

Increased nickel content for improved core toughness, slower response results in longer processing times.

Maximum nickel content for maximum core toughness, slower response results in longer processing times.

Table 29.23 REQUIRED CASE DEPTHS OF SELECTED APPLICATIONS FOR CARBURISED PARTS

Application ~t m

Case depth

inches

High wear resistance, low to moderate loading. Small and delicate machine parts subject to wear.

High wear resistance, moderate to heavy loading. Light industrial gearing.

High wear resistance, heavy loading, crushing loads or high magnitude alternating bending stresses. Heavv duty industrial gearing.

Bearing surfaces, mill gearing and rollers.

<508 <0.020

508-1 016 0.020-4).040

1 016-1 524 0.040-0.060

1 524-6 350 0.060-0.250

�9 Hardening can be performed after initially normalising the structure of the outer layer and core; this method is similar to double hardening, except that the steel is cooled slowly after carburising. The core possesses a normalised structure with less strength but more ductility.

Depending on the alloy selected and the component shape, the carburised steel may be quenched in water, oil or a water polymer solution. Specific recommendations are provided in SAE Standard AMS 2759.

Given the various limitations of both oils and water, there is continuing research into alternative quench media in the heat treating industry. For example, there has been a long search for quenchants, which would exhibit faster cooling rates than those produced by many quench oils, to avoid pearlitic microstructure formation in many steels. This has led to the development and use of aqueous polymer quenchants, as an alternative to many quench oils.

Water quenching of gas carburised parts will provide cooling rates in excess of the critical velocity for martensitic transformation. However, water quenching is likely to cause hardening cracks. A new quenching procedure has been developed using water as the medium and it is known as intensive quenching.

Quenching in oil, on the other hand, reduces the risk of cracking, but often it is not possible to achieve the desired martensitic microstructure throughout the entire carburised case. Only water soluble polymer quenchant solutions provide cooling rates ranging between those attainable with water and those achievable with oil. The Grossmann H-factor obtained for polymer solutions may vary from H = 0.2-1.2 (water typically exhibits an H-factor between 0.9 and 2, whereas the H-factors for oil may vary from 0.25 to 0.8). Therefore, the use of an aqueous polymer quenchant will provide a relatively 'mild' quench severity, sufficient to provide the desired martensitic transformation for both case and core. 56~1

After hardening the carburised case by quenching, it is necessary to temper the steel by selecting a low tempering temperature in the region of 180-275~C to reduce the residual stress caused by the phase transformation from austenite to martensite. A reduction in the quantity of retained austenite may be achieved by cryogenic treatment before tempering, especially after hardening variants '2 ' and '4 ' in Figure 29.23 in the case of alloy steels.

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%C

Treatment

1. Pot quenching (direct hardening).

2. Single hardening from a temperature optimised for the carburised surface.

3. Single hardening from an intermediate temperature.

4. Single hardening from a temperature optimised for the core.

5. Regenerative quenching (double hardening).

6. Pot quenching with isothermal holding.

Carburized surface

Dissolved carbides, retained austenite.

Excess of undissolved carbides.

Slightly coarsened grains, limited extent of carbide dissolution.

Slightly larger grains, well dissolved carbides, occurrence of retained austenite.

Limited extent of carbide dissolution, reduced occurrence of retained austenite.

Treatment with single stage diffusion of carbon, excess of carbides decays, reduced retained austenite content, minimum of deformation.

Core

Increased hardness.

Soft and workable.

Greater strength and more ductility than 1.

Peak hardened, with a better compromise between strength and ductility than 2.

Soft and workable, high level of ductility and resistance to high strain rate deformation (impulse strength).

Increased hardness.

Figure 29.23 Heat treatment after carburising (not drawn to exact scale)


Steels that are hardened after vacuum or ion carburising are often quenched in a high-pressure gas. The high-pressure gas may also be blended (helium/nitrogen) to achieve the desired cooling rates required for the alloy being quenched and the geometry of the part. The heat transfer coefficient of a gas quenching medium depends on the gas type, delivery pressure, velocity, turbulence and directionality. In addition to producing clean parts, a major advantage of gas quenching is that is does not harm the environment.

STRUCTURE-PROPERTY RELATIONSHIPS

Despite the progress and range of applications of carburising technology, the influence of case structure on properties continues to be controversial. One of the most significant issues is the influence, on properties, of retained austenite and carbides in the structure of case-hardened layers and of the carbon concentration in the surface zone of those layers. For example, an improvement in bending resistance has been observed for 10-30% retained austenite. However, others have reported an unfavourable influence without specifying the percentage values. ~'2 t,~ Reference 62 states that fatigue strength increases as the amount of retained austenite increases and recommends that the minimum retained austenite content should be about 25%. Other work has reported a detrimental influence without stating the bending resistance limit. < 6s

The differences of opinion on the influence of retained austenite are due to a lack of understanding of the detailed conditions and mechanisms involved in microstructure formation and consequently of the structure-property relationships of case-hardened layers. Variable process methodologies that have been utilised also contribute to the different results obtained. The problem is related to the ways in which retained austenite forms in the surface zone of carburised layers and the need to account for the influence, on properties, of changes in other layer parameters of case-hardened steels. These additional parameters include: the carbon content in hardened structures, chemical composition of the matrix, size of former austenite grains, martensite morphology, internal stress state, etc.

There have been reports that carbide precipitation does not unfavourably influence properties. 62"64-6v However, others have reported detrimental effects. These reports should be regarded with great caution, since lower properties than expected may also be due to an overabundance of carbon or tensile stresses in the surface zone of the layer, related to this overabundance. Also, the manner in which the carburised part is loaded can have an effect on the observed properties. The presence of carbides in the structure of carburised cases, with appropriate hardness, allows the production of machine parts and tools which will have high abrasive wear and bending resistance. The carbide form and fraction in the case structure will depend on how the carburising process is conducted, especially on the carbon potentials of the atmosphere, and the temperature and time of the process (see Table 29.25).

One of the factors that exhibits a great influence on the properties of machined parts and tools is grain size. Austenite grain formation in the diffusion layer and in the core influences the morphology of martensite being formed. 62,66,69.70 Austenite grain formation also affects the plastic properties of the carburised steel. Grain size variation during carburisation influences the resistance to bending fatigue 71 or the value of the fracture toughness KR.. "2 The small effect of austenite grain size in the case, or in the core, can be examined by controlling the formation of carbides, the chemical composition of the steel or through process control.

According to previous work ~3.74 austenitising from a two-phase austenite-cementite region reduces the occurrence of plate martensite microcracking. Austenitising in the temperature range where a two phase structure, austenite plus carbide is formed, leads to cracking of the martensite. Austenitising at this temperature is favoured for small austenite grain material. The favourable influence of austenitisation temperatures less than A,-m has been reported. ~5'6 The influence of austenitisation temperature on the properties exhibited by the carburised case has also been addressed 63"7v-v9 (see Table 29.25). In addition, the effect of the mechanical loading mode on the relationship between structure and properties of the diffusion layer has been considered 66 (see Table 29.25).

A low Ms temperature results in a reduction of the temperature range for which self-tempering of martensite may occur, which favourably affects many properties. In other words, a low Ms temperature leads to significant inhibition of self-tempering processes occurring during hardening, and, therefore, to a high carbon content in solid solution (in martensite) which will increase the strength and hardness. A minimum Ms value corresponds to the highest compressive stresses.

The presence of hard carbides can have a significant effect on the hardness of the overall coating. Carbide microhardness in carburised carbon steel has been measured at 1 000 HV, and in 2% Ni--Cr steel ~880 HV. 8~ Therefore, carburised and hardened surfaces containing 25 to 50% of dispersed

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carbides should exhibit higher hardness, 65 to 67 HRc (830 to 900 HV), s~ than carburised surfaces without carbides.

The presence of carbides in the case also has an indirect influence over the occurrence of internal stresses. The nature and extent of carbide precipitation's influence on the matrix depends on the quenching method. This influence can occur either as a result of the effect that carbide formation has on the chemical composition of the carbides' surroundings (a micro-scale effect), or the influence that the carbides themselves have on the overall stress state in the coating (a macro-scale effect). With slow cooling the micro-scale influence of the carbides on the matrix dominates, while for fast cooling the macro-scale influence is the important consideration. If the matrix of a case containing carbides transforms into martensite, the resulting macrostresses will probably be compressive. The magnitude of these stresses depends on the amount of retained austenite and on the martensite type (plate or lath) (see Table 29.25). Slow formation of large carbides will leave the eutectoidal carbon content in the matrix, which for high-alloy steels is a relatively low content. If the formation of increased amounts of plate martensite in the matrix is favoured during quenching, this results in higher compressive macrostresses.

It has also been shown ~2 that surfaces containing large amounts of carbides result in the formation of lower compressive stresses ( - 6 0 MPa) than a surface without carbides ( -500 MPa). Other work has shown how various carbides influence surface internal stresses (see Table 29.25). s3 In cases where the formation of carbides has led to lower compressive stresses, the extent to which this is desirable depends on the operating conditions under which the hardened carburised case is used.

Contact fatigue life for four different carburised surfaces is presented in reference [84], based on the results of sliding and rolling tests on carburised and tempered 2% Cr-Mn steel (see Table 29.25). Tests showed that coarse carbides can also (like coarse-grain martensite) influence contact fatigue (see Table 29.25).

The presence of sufficient amounts of hard carbides might be expected to guarantee that surfaces will have both abrasive and adhesive wear resistance. Indeed, it is generally found that increasing the surface carbide content will increase abrasive wear resistance, in cases where spheroidised carbides appear in a non-martensitic matrix. In contrast, when the structure of the matrix consists of martensite and retained austenite in various amounts, the amount of carbides present does not significantly influence wear resistance, st A microstructure that produces the best contact fatigue resistance also increases abrasive wear resistance, st' It should be further noted that contact fatigue and abrasive wear resistance are also influenced by the quality and condition of lubrication employed in the system.

Carbide network continuity, carbide coarsening and excessive penetration depth reduce fatigue resistance. A high manganese fraction in a low-alloy steel may be the cause of carbide networks in carburised cases. The form and amount of carbides are influenced not only by the diffusion process temperature and time, but also temperature and duration of subsequent treatments (e.g. annealing, or hardening of the system--see Table 29.25).

Table 29.25 provides factors relating structural and working properties of the carburised case. <

29.2.12 Carbonitriding

THE PHYSICAL-CHEMICAL BASIS OF THE CARBONITRIDING PROCESSES

Process overview

Carbonitriding is dependent on the simultaneous diffusion of carbon and nitrogen into the surface layers of the steel. The process is usually conducted in the temperature range of 850-880~C. The process gas for the treatment is based upon the use of either a nitrogen/methanol blended gas or an endothermic atmosphere, with the addition of a hydrocarbon enrichment gas as a source of carbon, plus ammonia to supply the nitrogen.

The carbonitriding process combines simultaneously nitriding and carburising, and while these two processes are distinguishable, they are not completely independent.

Deposition of carbon and nitrogen and diffusion into the steel

The diffusion of carbon from the process atmosphere into the steel surface is controlled by Fick's laws of diffusion (as discussed earlier in this chapter, with respect to carburising). However, the simultaneous diffusion of two interstitial solutes (i.e. carbon and nitrogen) in the austenite does raise the possibility of interactions between these, which might lead to a change in the diffusivity of each solute. Furthermore, both carbon and nitrogen are derived from the process atmosphere and undergo


adsorption onto the steel surface (followed by diffusion into the surface) and so the deposition reactions might act in concert. The above raises a question, 'what ability (if any) does nitrogen have to facilitate the carbon adsorption process and does the nitrogen assist carbon absorption into the steel surface?'

At the carbonitriding process temperature, the primary equilibrium thermal disassociation processes producing carbon and nitrogen are:

C+CO2 ~ 2CO Reaction l CH4 r C + 2H2 Reaction 1

C + H20 ~ CO + H2 Reaction 2 2NH3 r N2 + 3H2 Reaction 3

NH3 r N + 3/2H~ Reaction 4 1/2 N2 ~ N Reaction 5

HCN ~ C + N + 1/2H, Reaction 6 _

CO + 2NH3 r CH4 + H20 + N2 Reaction 7 CO + NH3 r HCN + H20 Reaction 8 CO2 + H2 r CO + H20 Reaction 9

CH4 + H20 ~ CO + 3H, Reaction 10 CH4 + CO2 ~ 2CO + 2H, Reaction 11

_

Reactions 1 to 6 play a direct and active part in the production of both carbon and nitrogen. In creating atomic carbon for the diffusion reactions, reactions 1-3 have the greatest significance. Atomic nitrogen is created in reactions 4-6 and is available to the steel for diffusion, together with the carbon. A key point is that, at the process temperature of the carbonitriding procedure, molecular nitrogen will decompose to create active, atomic nitrogen (N) in accordance with reaction 6. In accelerating the penetration of both carbon and nitrogen, part of the process will create very small quantities of prussic acid (hydrocyanic acid) on the steel surface. The remaining direct or indirect reactions are inconsequential.

Atomic hydrogen is emitted and is directed to the surface of the steel and will intensify both the reduction and dissociation of surface oxides, as well as converting the oxygen emitted as a result of the carburising reaction. When using 'real-world" furnace atmospheres, oxygen will always be present, as a result of moisture from the endothermic gas, or from unavoidable furnace leakages.

The source of nitrogen for diffusion into the steel is derived from the introduction of ammonia, together with the hydrocarbon enrichment gas. At the process temperature, ammonia dissociation will occur rapidly, which begins in reaction 4 and ends in the final decomposition described in reaction 5. Atomic nitrogen is produced almost immediately as result of the decomposition of the ammonia, followed by the decomposition of molecular into atomic nitrogen. Note that the volume of ammonia should be between 4% to 8% of the total gas flow into the process furnace.

The size of the N2 molecule is too large to permit this to diffuse into the austenite matrix of the steel and so nitrogen will not dissolve into the steel as molecular nitrogen (the solubility is limited to 0.025%). It is only the active or atomic nitrogen, resulting from the dissociated ammonia, that will penetrate the steel surface and diffuse into the austenite. The limit of solubility of atomic nitrogen in iron, at a temperature of 600~ is approximately six percent.

The quantity of dissociated ammonia, which is available for diffusion, is proportional to the un-dissociated ammonia present in the furnace atmosphere. It is not possible to quantify the un- dissociated ammonia on the basis of thermodynamic dependence, resulting from the conditions of equilibrium between ammonia and other products present in the furnace atmosphere. This is because such an equilibrium will only occur in the range of temperatures between 400 to 600~C. Between this range of temperatures, dissociation is almost total and the quantity of un-dissociated ammonia does not exceed 0.1%.

The factors influencing the degree of dissociation of ammonia, which occurs during the carbonitriding process are:

�9 Process temperature. �9 Atmosphere changes per hour within the furnace. �9 Circulation and distribution of the process gas within the furnace. �9 Furnace size (process chamber volume). �9 Furnace heating elements versus a gas heating system.

The surface properties of the carbonitrided layers are essential in determining the concentration of nitrogen in the superficial layer of the steel. This concentration will depend directly on that established during the early reaction stages of the ammonia (at the gas-metal interface) and indirectly on other


influencing factors such as the speed of the reaction. Therefore for the control of the process, the superficial concentration of nitrogen is usually a function of the concentration of ammonia in the atmosphere, for different temperatures.

At even higher temperatures of 850 to 930: C, the amount of ammonia will be increased, as a result of the temperature. This addition of ammonia will, to a small degree, increase the concentration of nitrogen in the surface layer of the steel. If this is allowed to occur, then the potential for the formation ofnitride networks is extremely high. Conversely, if the process temperature is reduced, the opposite will occur. Thus, even small additions of ammonia will cause a dramatic increase in the surface layer nitrogen concentration. Great care should be taken with the addition of ammonia, as a source of atomic nitrogen, as this can cause nitride networks to occur within the formed case.

Steels that include nitride forming elements (for example: aluminium, boron or silicon) will react readily with nitrogen. The alloying elements that will react favourably with nitrogen to form stable nitrides.

It should be noted that chromium, when combined with manganese, will display higher concentrations of nitrogen in the surface layers in comparison with steels that do not include these elements. Furnace atmospheres with higher carbon potentials can saturate the steel surface excessively. Con- sideration must be given not only to the gaseous reactions, but also to suppression, by the nitrogen activity at the steel surface, of the potential for retained austenite formation. Excess carbon will also encourage carbide formation in the surface of the steel, and nitrogen will assist in diminishing the retention of retained austenite, by allowing a lower austenitising temperature to be selected. An added benefit of the nitrogen reaction (particularly with the nitride forming alloying elements), is that there is an increase in the resulting surface hardness, as well as a reduction of the potential for distortion.

ROLE OF NITROGEN IN THE CARBONITRIDING PROCESS

The direct influence of carbon on the diffusivity of nitrogen in solid solution in the steel is comparatively insignificant. However. there are many other very essential influences of carbon on the process, which allow nitrogen to enter the steel. In the process of carbonitriding, a role of nitrogen is to facilitate the solutioning of carbon in iron. Nitrogen will also contribute to a significant increase in the speed of the carbon diffusion process.

Changes of carbon potential in the fio'nace atmosphere

Ammonia and the products of ammonia dissociation will change the equilibrium conditions of the main carburising reaction, thus continually changing the carbon potential of the furnace process atmosphere. The products of the dissociation of ammonia (nitrogen and hydrogen) are in accord with reaction (4) and this influences the partial pressures of the carburising components of the atmosphere, compared with non-carburising components and has the effect of raising the carbon potential of the furnace atmosphere. In contrast, ammonia will also react with oxides of carbon and will produce water vapour, which will lower the carbon potential of the furnace atmosphere. Thus, the completion of the ammonia dissociation reaction (and hence its ability to change the atmospheric carbon potential) is very important.

Ammonia will influence the dependence between the carbon potential of the endothermic atmosphere (and enrichment gas) and the temperature of the dew point. Seemingly insignificant additions of ammonia to a carburising atmosphere will cause a decrease in the carbon potential of the furnace atmosphere, which also means a reduction of the atmosphere's dew point.

Increasing the diffusion coefficient of carbon in austenite

The diffusion of nitrogen into steel raises the diffusion coefficient of carbon in austenite. This offers the possibility ofutilising low process temperatures. A carbonitride layer developed, at a given temperature, can have a thickness that is similar to that ofa carburised layer produced at a temperature that is approximately 50:C higher. Creation of the carbonitrided layer will begin below 850~C and the case will begin to form faster, than for carburising at a temperature of 900: C.

The activity of carbon in austenite

Nitrogen will increase the activity of carbon in austenite, as a result of the superficial concentration of carbon in the primary layer from the furnace atmosphere. This provides an additional reason why, in comparison with carburising, the process ofcarbonitriding can utilise a lower process temperature.


Changes in the equilibrium conditions in the Fe-Fe3C system

Changes in the equilibrium conditions in the iron-cementite system are caused by the presence of nitrogen. Nitrogen acts as a y stabiliser and reduces both the A1 and A3 temperatures. This assists the diffusion of carbon into the steel surface (the carbon solubility in austenite is much higher than that in ferrite) and gives the possibility of applying lower process temperatures for carbonitriding, as well as utilising lower temperatures for hardening.

Response to quenching

The introduction of nitrogen into the surface of the steel improves the ability to quench the steel without damage. Addition of nitrogen delays diffusional decomposition of the austenite and hence allows the use of a reduced cooling rate.

By the utilisation of lower austenitising temperatures, oil can be employed successfully as the quench medium, thus considerably reducing the potential for distortion. Because of nitrogen diffusion into the steel surface, the dimensional stability of the treated part is greatly improved.

Quenching conditions will determine the surface hardness, the type and properties of induced internal stresses and mechanical properties, such as torsional and tensile strength. Because of potential defects related to water quenching, it is generally recommended that carbonitrided steels be quenched in oil. The use of specially blended quench oils, that would facilitate martensitic transformation of the surface, whilst allowing a pearlitic core structure, within a specific range of oil cooling speeds is recommended. However, it is typically necessary to provide exhaust ventilation for gaseous waste products and there is often a concern for the environmental impact related to the use of quench oils.

It has been found that aqueous solutions of polyalkylene glycol (PAG) based quench media will reduce the potential environmental impact and by varying the concentration of the polymer quenchant in water, a wide range of cooling speeds is possible. The general impact of quenchants and cooling rates on the properties of carbonitrided steels will be discussed subsequently.

Reduction of the martensite start temperature

The presence of nitrogen in the carbonitrided layer will reduce the martensite start temperature (Ms), as well as significantly raising the martensite finish (Mr) temperature. Thus, the Ms temperature of the carbonitrided steel will be lower than if the same steel was carburised and will also cause a reduction in the compressive stress through the case.

Reduction of retained austenite

Lowering the martensite start temperature will reduce the amount of residual austenite in the carbonitrided layer. Residual austenite will normally be found when the surface carbon potential is in excess of 0.7% to 0.9%, with concentrations of nitrogen in solution. If the concentration of both carbon and nitrogen is controlled in the case, then the retention of any untransformed austenite will be minimal, thus producing a relatively dimensionally stable case.

Control of carbon and nitrogen within the diffusion layer

The production of available carbon for the process is derived principally from the enrichment gas, which is added to the endothermic gas, or the nitrogen/methanol plus enrichment gas. The objective is to ensure a surface carbon potential (0.7% to 0.9% carbon) around the eutectoid composition. The carbon potential can be controlled by accepted methods, such as: dew point, shim analysis, oxygen probe or CO/CO2 analysis. The nitrogen is normally derived from the ammonia enrichment gas added to the furnace. This is extremely difficult to control, as there is no commercial system that can be used to control accurately the decomposition of ammonia to release nascent nitrogen. Therefore one must rely on the relationship of the volumetric flow of the process gas (carrier gas and enrichment gas) to the ammonia. Generally, the ammonia flow should be between 4% to 8% (by volume) of the total gas flow into the furnace.

THE CARBONITRIDING PROCEDURE

The carbonitriding process very closely resembles the process of gas carburising. The procedure involves the following steps:

�9 Conditioning of the furnace. �9 Loading the work into the furnace.


�9 The carbonitriding process. �9 Quenching the carbonitrided work.

The carbonitriding process times are much shorter than soaking times experienced with the carburising process.

During the heating cycle of the batch, the introduction of the carbonitriding atmosphere into the furnace takes place. It is not necessary to introduce the process gas into the furnace during the heat up cycle. This is because during the heat up of the batch and particularly in the temperature range of 550~ to 700~ ammonia can cause the formation of nitrides at the steel surface. These nitrides may not be stable at the carbonitriding process temperature and may undergo transformation.

The selection of the carbonitriding process temperature requires careful consideration, with due regard to the atmospheric composition employed, as well as the heating method of the furnace. This is to ensure a well conducted process that is able to produce the required case depth, as well as the appropriate carbon and nitrogen concentrations within the case. When selecting either low carbonitriding process temperatures or long cycle times to produce deep cases, the probability of high surface nitrogen concentrations is extremely high.

It is not usual to select deep carbonitrided cases. Generally the case for carbonitriding has a maximum depth of 500~tm (0.020in.). Deep case formation will give rise to a high probability of the retention of un-transformed austenite (retained austenite). If ammonia is introduced into the furnace, for example during the last one-third of the process cycle, then the risk of high surface nitrogen concentrations is reduced almost to the point of elimination. On completion of holding of the workload in the furnace atmosphere, having allowed sufficient time at the process temperature to form the desired case, the load is quenched from the appropriate austenitising temperature. Care must be taken when selecting the austenitising temperature, due to the potential risk of retained austenite and of course the potential for the occurrence of distortion.

HEAT TREATMENT AFTER CARBONITRIDING

Once the process of carbonitriding is completed, this is usually followed by a cooling procedure, down to the hardening temperature. A light tempering process at 180-200~ usually follows this. This is intended to temper the martensitic case and reduce the potential for case/surface cracking.

The cool-down procedure is selected to accomplish the formation of the structures that are needed to produce the properties required by the engineering design for the core and surface of the steel. The steel can be:

�9 Cooled within the furnace down to the case austenitising temperature in a controlled manner, followed by quenching.

�9 Cooled down to, perhaps, 500~ removed from the furnace and cooled externally, followed by reheating to the austenitising temperature and then by a quenching procedure which will harden the carbonitrided case.

Examples of heat treatments after carbonitriding may be found in Figure 29.24.

THE INFLUENCE OF STRUCTURAL FACTORS ON THE PROPERTIES OF CARBONITRIDED LAYERS

Table 29.26 provides information on the relationship between structural factors and the properties of carbonitrided materials. 62

Table 29.26 The influence of structural factors on the properties of carbonitrided layer. 77"1~176

STEELS USED IN CARBONITRIDING PROCESSES

Some steels used internationally for carbonitriding may be found in Table 29.27.


Treatment

1. Pot quenching (direct quenching).

2. Pot quenching with isothermal holding.

Carbonitrided surface

Dissolved carbides, retained austenite.

Treatment with single diffusion of carbon and nitrogen, excess of carbides or carbonitrides decays, reduced content of retained austenite, minimum of deformation.

Core

Increased hardness.

Increased hardness.

Figure 29.24 Heat treatment a#er carh~mitridin~ (m~t drawn m exact scale)

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Table 29.27 THE CHEMICAL COMPOSITION OF STEELS USED INTERNATIONALLY IN THE CARBONITRIDING PROCESSES

.4verage composition (% )

Grade C M n N i Cr M o Coun tin"

20MoCr4 0.20 0.70 0.40 0.45 25MoCr4 0.25 0.70 0.40 0.45 20CrMo2 0.20 0.70 0.60 0.35 Germany 4028 0.24 0.82 0.25 20NiMoCr6 0.20 0.70 1.60 0.50 0.45 23CrMoB3 0.23 0.80 0.70 0.35 + B 8620 0.20 0.80 0.55 0.50 0.20 4024/28 0.23/0.28 0.80 - - 0.25 Great Britain Mod8822 0.22 0.85 0.55 0.50 0.55./0.65 18Co4 0.18 0.75 1.05 0.22 France 8620 0.20 0.80 0.55 0.50 0.20 8620 0.20 0.80 0.55 0.50 0.20 Italy 18NiCrMo5 0.18 0.70 1.35 0.80 0.20

29.3 Hea t t r e a t m e n t o f a l u m i n i u m alloys

In this section, a brief review of aluminium alloy physical metallurgy, as affected by heat treating, is provided. Defects arising from heat treating operations are also discussed. The importance of selecting the proper quenchant, and quenching parameters are explained.

29.3.1 Introduction to aluminium heat treating

Aluminium in many forms has been used in aircraft since their very beginning. This is because aluminium alloys can be heat-treated to relatively high strengths, while maintaining low weight. These alloys are easy to bend and machine, and the cost of the material is relatively low. Because of these advantages, aluminium is the most common material used in aerospace today. It is used in the manufacture of advanced commercial aircraft such as the Boeing 777, Airbus 380, and military aircraft such as the Boeing UCAV or F/A-18 E/E

Typically, the aluminium alloys used in aerospace structures are the heat treatable grades, such as 2XXX, 6XXX and 7XXX. 7XXX alloy grades such as 7075, 7040 and 7050 predominate in aerospace structures. Some of the A1-Li alloys such as 2090 and 8090 have also found application. These alloys are used commonly because they are easily available and readily formed and heat-treated to yield high strength, corrosion resistant components.

However, the use of aluminium is not limited to the aerospace industry. It is used extensively in, for example, sporting equipment and in the automotive industry.

The aluminium alloys most commonly used are 7XXX wrought products. The thickness of parts used ranges from 0.6 mm to 250 mm. Aluminium alloy 7075 and 7050 extrusions are used extensively for stringers and other structural requirements. Some extrusions of 2024 and 2014 are also used. Sheet, both clad and non-clad, of 2024, 7475 and 7075 is used for wing and fuselage skins. Sheet is also formed and built-up to produce bulkheads and other structural requirements. For heavy loading, large forgings of 7050 and 7040 are used commonly, particularly in military aircraft.

Recently, there has been much interest in the use of heat treated 7050 and 7040 plate to minimise the cost of forgings. In this case, the plate (heat treated at the aluminium mill) is supplied to the airframe manufacturer. There the plate is heavily machined to fabricate large ribs and bulkheads. The advantage of this is that it avoids heat-treating small parts and may serve to improve distortion and residual stress control. Unfortunately, there is a significant property variation through the thickness of thick plate. This variation has been studied in detail and the microstructure development in thick sections has been researched extensively, l~

Primarily, because of affordability, castings have also been used. However, the use of castings has been limited because of design factors and their limited ductility. Typically casting alloys such as A356 and A357 are used for casting. Applications of castings are simple, non-flight critical components, such as door handles and avionics cabinets.

Heat treatment o f aluminium alloys 29--67

Heat-treating of aluminium requires stringent controls. To achieve repeatable results, and to provide a quality product, the airframe manufacturers, suppliers and heat-treaters have developed a series of specifications. The most widely used specification is AMS-2770 'Heat Treatment of Aluminum Parts'. 1~ This specification details solution treating times, temperatures, quenchants and ageing practices. It also defines the required documentation for heat treat lot traceability. Furthermore, this specification establishes the quality assurance provisions needed to ensure that a quality product is provided to the airframe manufacturer.

29.3.2 A brief description of aluminium physical metallurgy

There are four aluminium alloy types that dominate the heat treatable alloys used in industry today. These are: AI--C u, A1--C u-M g, AI-M g-Si and A I-Z n-M g-C u. A brief ex p lanation o f the prec ipi tati on sequence, for each alloy type follows.

29.3.2.1 AI-Cu alloys

The aluminium-copper system has been reviewed in detail. 1 lo The equilibrium phase diagram TM

contains an eutectic, at 548~ and approximately 32% copper, involving the face centred cubic A1 solid solution phase (ct) in equilibrium with CuAI2 (0). The extent of solid solubility at the aluminium rich end is approximately 5.7% copper. Commercial alloys of this type are 2219, 2011 and 2025.

The precipitation sequence was originally established by Guinier 1121~3 and Preston. 114~5 Hornbogen further examined the precipitation in AI-Cu and confirmed the results of Guinier and Preston.116 The precipitation sequence, on ageing after rapid quenching, has been accepted as being Guinier-Preston zones (GPZ) in the form of plates parallel {001 }.~,l, transforming to the coherent precipitate 0", followed by semi-coherent 0" plates parallel { 001 }.~,l. The final equilibrium precipitate is 0(Cu2A1). Silcock et al., 11"~ examined this progression of precipitates and showed multiple stages in precipitation, evidenced by changes in hardness and Laue reflections.

29.3.2.2 AI-Cu-Mg alloys

Aluminium-copper-magnesium alloys were the first precipitation hardenable alloys discovered, l ~s The first precipitation hardenable alloy was a precursor to alloy 2017 (4% Cu, 0.6% Mg and 0.7% Mn). A very popular alloy in this group is 2024.

The addition of magnesium greatly accelerates precipitation reactions. In general, the precipitation sequence, starting from a supersaturated solid-solution (SSSS), is:

SSSS ~ GPzones - -~ S'(AIzCuMg) ~ S(AI2CuMg)

The GP zones in this system are generally considered to be collections of Cu and Mg atoms collected as disks on the {ll0}~l planes. S' is incoherent and can be observed directly in the transmission electron microscope (TEM). S' precipitates heterogeneously on dislocations. These precipitates appear as laths on the { 210} ~l, oriented in the <001> direction. 11~ Since S' precipitates on dislocations, cold working after quenching increases the number density of S' and produces a fine distribution of precipitates in the matrix.

29.3.2.3 AI-Mg-Si alloys

The aluminium-magnesium-silicon alloy system forms the basis for the 6XXX series aluminium alloys. In this heat treatable alloy system, magnesium is generally in the range of 0.6-1.2% Mg and silicon is in the range of 0.4-1.3% Si. The sequence of precipitation is the formation of GP zones, followed by metastable/3' (Mg_~Si), followed by the equilibrium fl (Mg2Si). The GP zones in this case, are needles oriented in the <001> direction, with/3' and/3 showing similar orientations.

29.3.2.4 AI-Zn-Mg-Cu alloys

This aluminium-zinc-magnesium-copper series of alloys is probably the most popular and readily used. These alloys are used extensively in the aerospace industry because they have excellent corrosion resistance, fracture toughness and strength, compared with the other age hardenable aluminium

29---68 H e a t t r e a t m e n t

alloys. In 7XXX A1-Zn-Mg-Cu alloys, several phases have been identified that occur in A1-Zn-Mg- Cu alloys as a function of precipitation sequence. Four precipitation sequences have been identified. These are shown schematically below:

C~s~ss =:~ S

~ssss =:~ T' :=~ T

C~ss~s =~ VRC =~ GPZ ~ r/' =~ r/

C~ssss :=~ r/

In the first precipitation sequence, the S phase, A12CuMg, is precipitated directly from the supersaturated c~ solid solution (O~ssss). The S phase is reported to be orthorhombic, ~2~ with a space group of Cmcm, and 16 atoms per unit cell. The lattice parameters are: ~2~ a=401 pm, b = 925pm, and c = 715 pm. This phase has been identified ~:: as a coarse intermetallic that is insoluble in typical A1- Zn-Mg-Cu alloys at 465~ and as a fine lath precipitate in AI-4.5%Zn-2.7%Cu-2.2%Mg-0.2%Zr alloys. No orientation relationships to the matrix have been identified in the literature.

In the second precipitation sequence, an intermediate phase T' occurs in the decomposition of the supersaturated solid solution. Bernole and Graf first identified this phase. 123 Auld and McCousland 124 suggested that the structure was hexagonal with the reported lattice parameters a = 1.39nm and c = 2.75 nm. It was further suggested that the orientation of the hexagonal cell to the aluminium matrix is:

(0001)T'//(111)A1 (10i0)T'//(112-)A1

Further on, in the second precipitation sequence, the equilibrium T phase forms. This phase was reported ~25 to be cubic, space group Im3, with 162 atoms in the unit cell. It was indicated that the lattice parameter varies from 1.41 to 1.47 nm, with this variation being due to compositional variations. Bergman et al. 126 have proposed that the chemical formula Mg32(A1,Zn)49 is appropriate. The T phase was found to be incoherent with the aluminium matrix. This phase has been rarely reported in substantial quantities, even though commercial heat treatments up to 150~ lie in the A1 + MgZn2 + Mg32(A1,Zn)4~ phase field. In general, the T phase only precipitates above 200~ 12s

In the third sequence of precipitation, the supersaturated solid solution decomposes to form vacancy-rich clusters, Guinier-Preston zones, r]' and then 1/. Guinier-Preston zones have been inferred in A1-Zn-Mg alloys, based on small increases in electrical conductivity and an increase in hardness during the initial stages of ageing. ~_~4

The 7/' phase is an intermediate step towards the precipitation of the equilibrium phase r/(MgZn2). Direct evidence of 7/' is rare, and difficult to obtain. It has recently been accepted that the 1/' phase is hexagonal, however, the reported lattice parameters vary widely.

29.3.3 Defects associated with heat treatment

During the production of a part, defects can occur. These can come from operations before heat treatment, such as midline porosity or inclusions that are formed during casting the ingot. Further defects can form during homogenisation of the ingot, such as segregation, the formation of hard intermetallics and other second phase particles.

Defects associated with the heat treatment of aluminium can occur during solution treatment, quenching or ageing. Solution treatment defects include oxidation, incipient melting and under- heating. Defects that occur during quenching are typically distortion of the part or inadequate properties caused by a slow quench, resulting in precipitation during quenching and inadequate supersaturation. Defects that can occur during ageing include growth or shrinkage of the part. Problems can also arise from an inadequate response, by the material, to the ageing treatment. Heat-treatment-related defects are discussed in more detail below, in the sections dealing with the relevant process step.

29.3.4 Solution treatment

Aluminium alloys are classified as either heat treatable or not heat treatable, depending on whether the alloy responds to precipitation hardening. In the heat treatable alloy systems, such as 7XXX,

Heat treatment o f aluminium alloys

Table 29.28 SOLUTION TREATMENT TE\IPERATLRE RAN(iE AND EUTECTIC MELTING TEMPERATURE FOR 2XXX .-\LLOYS

Solution treatment Eutectic' melting Allov temperature range (: C) temperature (- C)

2014 496--507 51 (I 2017 496-507 513 2024 488-507 502

29-69

6XXX and 2XXX, the alloying elements show greater solubility at elevated temperatures than at room temperature. This is illustrated for the A1-Cu phase diagram (see Chapter 11 ). ~-~"

The AI-Cu phase diagram shows that holding a 4.5% Cu alloy at 515 to 550 ~ will cause all the copper to go into solution completely. The temperature used to achieve complete solutioning is known typically as the 'solution heat treating temperature" or 'solution treatment temperature'. If the alloy is cooled slowly, then the equilibrium structure of a (saturated solid solution of Cu in AI) + CuA12 will form. The CuAlz that forms is large, coarse and incoherent. However, if the alloy is cooled rapidly, there will be inadequate time for the CuAlz to precipitate. Hence, in the rapidly cooled alloy, all the solute is held in a supersaturated condition. Controlled precipitation of the solute, as finely dispersed particles, at room temperature (natural ageing) or at elevated temperatures (artificial ageing) is used to develop the optimised mechanical and corrosion properties of these alloys.

Solution treatment involves heating the aluminium alloy to a temperature slightly below the eutectic melting temperature. Solution treatment develops the maximum amount of solute into solid solution. This requires heating the material close to the eutectic temperature and holding at temperature long enough to allow close to complete solid solutioning. After solution treatment, the material is quenched to maintain the solute in supersaturated solid solution.

Because the solution treatment temperature is so close to the eutectic melting temperature, temperature control is critical. This is especially true for 2XXX series alloys. In this alloy group, the eutectic melting temperature is only a few degrees centigrade above the maximum recommended solution treatment temperature (Table 29.28).

29.3.4.1 Oxidation

If parts are exposed to temperature for too long a time, what is commonly referred to in the industry as 'high-temperature oxidation' could become a problem. Note: the term high-temperature oxidation is really a misnomer, in the context of the heat-treatment of aluminium alloys. Instead, in the present case, the culprit is actually moisture present in the air, during solution treatment. This moisture is a source of hydrogen, which diffuses into the base metal. Voids form at inclusions or other discontinuities. The hydrogen gas gathers and forms a surface blister on the part. In general, 7XXX alloys are the most susceptible (particularly 7050), followed by the 2XXX alloys. Extrusions are the most prone to blistering, followed by forgings.

Eliminating the moisture will minimise the problem of surface blistering. This is accomplished by the sequencing of doors over quench tanks and thoroughly drying and cleaning furnace loads prior to solution treatment. It is also important to make sure that the load racks used for solution treatment are dry. However, it is not always possible to eliminate high humidity in the air, in order to prevent surface blistering. Often the ambient relative humidity is high, so that other measures may have to be taken.

Ammonium fluoroborate is typically used to prevent blistering on 7XXX extrusions and forgings. An amount equivalent to 5 g per m 3 of workload space is usually employed to prevent surface blistering. This is applied as a powder, in a shallow pan, hanging from the furnace load rack. This material is very corrosive and requires operators to wear the appropriate personal protective safety equipment. Because the material is corrosive at temperature, it is recommended that the inside panels in the furnace be manufactured from stainless steel. The use of stainless steel panels will reduce corrosion and hence maintenance.

Anodising parts prior to solution treatment is an alternative to ammonium fluoroborate. This is generally practical for larger extrusions and forgings, where the cost of anodising is small compared with the cost of the part. However, for small parts, the additional cost does not generally justify the possible benefit of anodising prior to solution treatment.

29-70 Heat t reatment

29.3.4.2 Eutectic melting

Non-equilibrium conditions can occur because of localised solute concentrations. Because of the increased concentration of solute, the eutectic temperature could be decreased, causing localised melting. This is often called incipient melting. When this occurs, significant decreases in properties result. Properties most affected include toughness and tensile properties (both strength and ductility).

Local melting can also occur if the material is heated too quickly. This is particularly true of 2XXX alloys. In this alloy system, there are local concentrations of Al2Cu. At slow heating rates, the AlzCu dissolves slowly into the matrix. At high heating rates, there is inadequate time for the AIzCu to dissolve. Local concentrations cause the local eutectic temperature to drop, resulting in localised melting. If inadequate time is allowed for this metastable liquid to dissolve into the matrix, then in general, there is no decrement in properties.

Based on the equilibrium solidus temperature, 7XXX series aluminium alloys should be safe from eutectic melting. However, these alloys can exhibit eutectic melting under certain circumstances. In the 7XXX series alloys, there are two soluble phases, MgZn2 and A12CuMg. AI2CuMg is very slow to dissolve during solution treatment. Local concentrations of AI2CuMg can cause non-equilibrium melting, between the temperatures of 485 and 490~ (905 to 910~ and this is a problem, since the work may be brought to this temperature range, or a higher temperature. ~2s Homogenisation practice of the ingot is the primary source of S phase in these alloys. Because of the hazards of eutectic melting, it is imperative that the 7XXX alloys be homogenised.

29.3.4.3 Under-heating

Under-heating during solution treatment can cause problems, by not allowing enough solute to go into solid solution. This means that less solute is available during subsequent precipitation hardening reactions. As an illustration of this, Figure 29.25 shows the effect of solution-treating temperature on the yield strength and ultimate tensile strength. As the temperature is increased for both alloys shown in the figure, the tensile strength is also increased. For 2024-T4, it can be seen that there is a change in slope and a rapid rise in final properties as the solutionising temperature is increased past about 488~

29.3.4.4 Furnaces f o r solution treatment

Furnaces used for the solution treatment of aluminium alloys are typically of two types. Either 'drop- bottom' furnaces or salt bath furnaces are used. Both are batch type operations. There are several continuous types of aluminium furnaces, but these are typically limited to smaller parts.

500

450

400

.ff c"

~ 350

300

/o 0- j

o

f ~ j f - f

i

-e-- 6061-T6 Sy ~ 6061-T6 Su

-k- 2024-T4 Sy -o- 2024-T4 Su

~ J

j ~

7

250 ....... 470 480 490 500 510 520 530

Solution heat treatment temperature

Figure 29.25 Tensile strength as a fhnction of solution-treating temperature /br 6061 and 2024 (after Hatch) 12s

Heat treatment o f aluminium alloys 29-71

Solution treating and quenching of these alloys is typically accomplished in large high-temperature ovens. In some applications, the furnace is supported above the quench tank. The quench tank moves under the furnace on rails. Sometimes there is more than one quench tank, with each tank containing a different quenchant.

Drop bottom furnaces are typically arranged with a moveable quench tank underneath. Often, there is more than one quench tank, one for water and the other for polymer quenchants. These quench tanks move back and forth as desired, and the selected quench tank is positioned under the furnace. The door on the bottom of the furnace opens and the workload is immersed into the quenchant. Sometimes, because the quench tank is so large, it is impractical to move the tank. Therefore, the furnace is moved in a similar fashion to a gantry furnace.

29.3.5 Quenching

In this section, the effects of quenching on the structure and properties of aluminium alloys are discussed. Further detail on quenching media may be found in the Appendix (Section 29.3.7).

The key consideration during quenching is to prevent precipitation, without damaging the part. The problem of undesired precipitation during quenching can be understood in terms of the kinetics of nucleation and growth, during diffusional phase transformations. 12s As with other diffusional phase transformations, the kinetics of precipitation occurring during quenching are dependent on the degree of solute supersaturation, versus the rate of solute diffusion, both of which are functions of temperature. The degree of supersaturation increases with undercooling and hence with decreasing temperature. Hence the driving force for nucleation increases as the temperature decreases. Conversely, solute diffusion coefficients, and hence the rate of diffusion-controlled growth of precipitates, increase with increasing temperature. When either the solute supersaturation or diffusivity is low, the overall rate of precipitation is low. At intermediate temperatures, the amount of supersaturation is relatively high, as is the rate of diffusion. Therefore the overall rate of precipitation is greatest at intermediate temperatures. This is shown schematically in Figure 29.26. The amount of time spent in this critical temperature range is governed by the quench rate.

Precipitation occurring during quenching reduces the amount of subsequent hardening possible. This occurs because solute is precipitated from solution during quenching, is unavailable for any further precipitation reactions. Precipitation during quenching results in lower as-aged tensile strength, yield strength, ductility and fracture toughness.

Quantifying quenching and the cooling effect of different quenchants, has been studied extensively. ~29-~32 The first systematic attempt to correlate properties to the quench rate in AI- Zn-Mg--Cu alloys was performed by Fink and Wiley 133 for thin, 1.6ram (0.064") sheet. A time-temperature-tensile property curve was created and was probably the first instance of a TTT- type diagram for aluminium alloys. It was determined that the critical temperature range for 75S is 400~ to 290~ This is similar to the critical temperature range found for AI-Zn-Mg-Cu alloys. 134

E I.-

1 ~ - . i / \

f

f

/ z .... /

\

\

\ _

f/ ,,

Rate

Precipitat ion - - - Supersaturat ion Diffusion rate

Figure 29.26 Schematic shouing the interrelationship hetween the amount o/solute supersatm'ation and diffhsion rate on the amount o/heterogeneoltS precipitation occurring during quenching


At quench rates exceeding 450~ s -~ . it was determined that maximum strength and corrosion resistance were obtained. At intermediate quench rates, from 450 down to 100 ~ C s- ~, the strength obtained was lowered (using the same ageing treatment), but the corrosion resistance was unaffected. Between 100 and 20~ s -~ , the strength decreased rapidly, and the corrosion resistance was at a minimum. At quench rates below 20~ s -~ , the strength decreased rapidly, but the corrosion resistance improved. However, for a given quenching medium, the cooling rate through the critical temperature range was invariant no matter the solution treatment temperature.

One method that quantifies the quench path and the material kinetic properties is called the 'quench factor' and was originally described by Evancho and Staley. ~3~ This method is based on the integration of the area between the time-temperature-property curve and the quench path. Wierszykkowski ~-~' provided an alternative explanation of the underlying principles of the quench factor. However, his discussion is more generally applied to the thermal path prior to isothermal transformation. The procedures for developing the quench factor have been well documented. 137 This procedure could be used to predict tensile properties, 143 hardness TM and conductivity. ~3'~ It was found that the ~uench factor could not be used to predict elongation because of its strong dependence on grain size ~:~'. This method tends to overestimate the loss of toughness. ~4z The quench factor method also can be used to determine the critical quench rate for property degradation. ~45

29.3.5.1 Quench rate effects

The objective of quenching is to preserve the solid solution formed during solution-treating. This is accomplished by cooling rapidly to room temperature. This rapid cooling maintains the vacancy concentration necessary to enable the low temperature diffusion required for GP zone formation. Solute that precipitates, at either the grain boundaries, or dispersoids, is lost and cannot be utilised for further strengthening. Hence, the best properties (with respect to both strength and corrosion resistance) are usually achieved by employing the highest possible quench rate.

As an illustration of quench rate effects in a 7050 alloy, a Jominy end quench bar was fabricated from extruded 7050, and quenched using the Jominy apparatus, l~ Quench rates ranging from in excess of 1 000~ s -I to 1-2~ s -~ were obtained. This bar was sectioned and examined at specific locations along the quenched bar, using a TEM. Distinct differences in the microstructure were observed. At fast quench rates, the GP zones are small, and exhibit strain fields around the zones. There is little evidence of precipitation at the grain boundaries. As the quench rate is decreased, the strain fields are still present, however they are not as intense. There is also evidence of precipitation on the grain boundaries, although these precipitates are small. Some precipitation also occurs within the grains. Again, these precipitates are small. As the quench rate is further decreased, larger and coarser precipitates are evident. These precipitates occur at the grain boundaries and in the interior of the grains. The lack of a strain field around the precipitates indicates that the precipitates are incoherent with the matrix. The solute contained in the coarse precipitates at the grain boundaries and in the interior of the grain are unrecoverable, and will not participate in any strengthening of the alloy. Further, the presence of the precipitates at the grain boundaries can serve as localised galvanic cells, providing a preferred path for intergranular and stress-corrosion-cracking.

29.3.5.2 Warpage and distortion

Of all the possible 'defects' occurring during the heat treatment of aluminium, distortion during quenching is the most common. It is probably responsible for most of the non-value-added work (straightening) and costs associated with aluminium heat-treating.

Distortion during quenching is caused by differential cooling, and differential thermal strains developed during quenching. These thermal strains could be developed centre-to-surface, or surface- to-surface. This differential cooling can be caused by large quench rates, so that the centre is cooled much slower than the surface (non-Newtonian cooling) or by non-uniform heat transfer across the surface of the part.

Aluminium alloys are more prone to quenching distortion than steels. This is because solution treatment temperatures are so close to the liquidus temperature in the former case. Aluminium alloys exhibit less strength and greater plasticity than steels at the solution treatment temperature (or austenitising temperature for steels). Furthermore, much higher quench rates are necessary in aluminium alloys, than in steels, to prevent premature precipitation from occurring during quenching and to maintain supersaturation of the solute.

In steels, there is a phase transformation from austenite to martensite, on quenching. This causes around a 3% volume change, during quenching. There is no analogous phase transformation in

Heat treatment of aluminium alloys 29-73

aluminium alloys that can cause cracking or distortion. However, the coefficient of linear expansion ofaluminium is approximately twice that of steel (2.38 x 10- 5 K- ~ compared with 1.12 x 10- 5 K- ~ ). This causes much greater changes in length or volume as a function of temperature, and increases the probability that distortion will occur.

Proper racking of the parts is critical. The parts should be fully supported, with the loads spread out over a large area, since the creep strength of aluminium is poor. Parts should be wired loosely to prevent the parts from hitting each other during solution treatment. If wired too tightly, the wire could cut the parts. The use of pure aluminium wire minimises this problem.

Because of the poor strength of the solution-treated aluminium parts, distortion of the parts can occur as they enter the quenchant. As a general rule. parts should enter the quenchant aerodynamically to avoid distortion of the part, before it even enters the quenchant. The parts should enter the quenchant smoothly--they should not 'slap" the quenchant.

Racking a part so that it enters the quenchant smoothly also offers the benefit that it is more likely to have uniform heat transfer across the part. Distortion is more likely to occur because of horizontal changes in heat transfer than by vertical differences in heat transfer.

29.3.5.3 Grain boundary precipitation

Grain boundary precipitation and the formation of a precipitate free zone (PFZ), are the result of solute coming out of solution during quenching. With a perfect quench, all solute would be held in supersaturated solid solution and no precipitation would occur during quenching. However, a perfect quench is rare with real parts of the types used typically in contemporary practice. Since there is little likelihood of a perfect quench, some precipitation almost invariably occurs. The degree of precipitation depends on the alloy and the quench rate.

As already discussed, the rate of precipitation during quenching is based on two competing factors: supersaturation and diffusion. As temperature is decreased during quenching, the amount of supersaturation increases, providing increased driving force for precipitation. In contrast, at the beginning of quenching, the temperature is high. increasing the rate of diffusion. The Avrami precipitation kinetics for continuous cooling can be described by: l~~

~" = 1 - exp(kr) n (1)

where ~" is the fraction transformed, k is a constant and r is defined as:

f dt (2) r - C---~

where r is the quench factor, t is the time (s) and Ct is the critical time (s). The collection of the Ct points for continuous cooling, also known as the C-curve, is analogous to the time-temperature- transformation curve for isothermal transformation. These equations can be used to predict the volume fraction that has precipitated, and how that precipitation will affect the properties.

To avoid excessive precipitation during quenching, three requirements must be met. First, the transfer time from the solution treatment furnace into the quench tank must be minimised. Second, the properties of the quenchant must be selected to enable a quench that is fast enough to ensure that proper supersaturation is achieved, so that the desired properties can be obtained during subsequent ageing. Lastly, the quench tank must have adequate thermal inertia so that the quenchant does not heat excessively, causing an interrupted quench. In addition, the quenching system must extract heat uniformly to minimise property variations.

The quench delay time, or the transfer time from the furnace to the quench tank, is defined differently for air furnaces and salt baths. For air furnaces, the quench delay time is defined as the time interval, from when the furnace door first begins to open, until the last corner of the workload is immersed into the quench tank. For salt baths, the quench delay time is defined as the time interval from when the first corner of the workbasket is exposed, to the time at which the last corner of the workload is immersed into the quench tank. In general, this time interval is independent of the alloy, but depends on the solution treatment temperature, the velocity of movement and the emmisivity of the workload. Table 29.29 shows typical allowable quench delay times for various thicknesses.

The specification for the maximum allowable quench delay time is based on the assumed amount of cooling of the workload, before it enters the quenchant. In general, the maximum allowable quench delay times can be exceeded if it can be demonstrated that the part temperatures do not fall below approximately 413~ before immersion. An exception to this is for AA2219, where the part temperatures can not fall below 482=C before immersion.


Table 29.29 TYPICAL MAXIMUM QUENCH DELAY TIMES

Minimum thickness

mm inch Maximum time (s)

Up to 0.41 Up to 0.016 5 Over 0.41-0.79 Over 0.016---0.031 7 Over 0.79-2.29 Over 0.031-0.090 10 Over 2.29 Over 0.090 15

It is difficult to measure directly and control the temperature drop during transfer of the workload from the solution treating furnace to the quench tank. However, the quench delay time is easily controlled using only a stopwatch. This is augmented with the results from routine tensile testing and intergranular corrosion testing.

29.3.6 Ageing (natural and artificial)

29.3.6.1 Natural ageing

Some heat treatable alloys, especially 2XXX alloys, harden appreciably at room temperature to produce the useful tempers T3 and T4. These alloys, when naturally aged to the T3 or T4 tempers, exhibit high ratios of ultimate tensile strength/yield strength and also have excellent fatigue and fracture toughness properties.

Natural ageing, and the resulting increase in properties, occurs by the rapid formation of GP (Guinier-Preston) zones from the supersaturated solid solution, by means of solute diffusion involving quenched-in vacancies. Strength increases rapidly, with properties becoming stable after approximately 4-5 days. The T3 and T4 tempers are based on natural ageing for 4 days. For 2XXX alloys, improvements in properties after 4-5 days are relatively minor, and become stable after one week.

The A1-Zn-Mg-Cu and A1-Mg-Cu alloys (7XXX and 6XXX), harden by the same mechanism of GP zone formation. However, the properties resulting from natural ageing are less stable. These alloys still exhibit significant changes in properties, even after many years.

The natural ageing characteristics vary from alloy to alloy. The most notable differences are the initial incubation time for changes in properties to be observed, and the subsequent rate of change in properties. Ageing effects are suppressed with lower than ambient temperatures. In many alloys, such as 7XXX alloys, natural ageing can be nearly completely suppressed by holding at -40~

Because of the very ductile and formable nature of as-quenched alloys, retarding natural ageing, increases scheduling flexibility for forming and straightening operations. It also allows for uniformity of properties during the forming process. This contributes to a quality part. However, refrigeration at the temperatures employed normally, does not completely suppress natural ageing. Some precipitation still occurs. Table 29.30 shows typical temperature and time limits for refrigeration.

29.3.6.2 Artificial ageing

After quenching, most aerospace aluminium alloys are aged artificially. This is a complex process and requires an understanding of vacancies, and the interaction of precipitation and metastable phases. In general, the sequence of precipitation occurs by clustering of vacancies, formation of GP zones, nucleation of a coherent precipitate, precipitation of an incoherent precipitate and finally the coarsening of the precipitates.

Precipitation hardening is the mechanism where the hardness, yield strength and ultimate strength dramatically increase with time at a constant temperature (the ageing temperature) after rapidly cooling from a much higher temperature (the solution treatment temperature). This rapid cooling or quenching results in a supersaturated solid solution and provides the driving force for precipitation on ageing. The age hardening phenomenon was first discovered by Wilm, ~ ~8 who found that the hardness of aluminium alloys with small quantities of copper, magnesium, silicon, and iron, increased with time, after quenching from a temperature just below the melting temperature.

H e a t t r e a t m e n t o f a l u m i n i u m a l l o y s

Table 29.30 TYPICAL TIME AND TEMPERATURE LIMITS FOR REFRIGERATED PARTS STORED IN THE AS-QUENCHED (AQ} CONDITION

29-75

Alloy

.~laximum storage time.lbr retention o f the.4Q condition (days)

Max imum delay time - 12~'C - 18:C -23:C after quenching ( 10 ~ F ) ( 0 ~ F } { - 10 ~ F ) (min) Max. Max. Max.

2014 15 1 30 90 2024 2219

6061 30 7 30 90 7075

Precipitation hardening (ageing) involves heating the alloyed aluminium to a temperature in the 95 to 230~ (200-450~ range. In this range, the supersaturated solid solution, created by quenching from the solution treatment temperature, begins to decompose. Initially there is a clustering of the solute atoms. Once sufficient solute atoms have diffused to these initial clusters, coherent precipitates form. Because the clusters of solute atoms have a lattice parameter mismatch with the aluminium matrix, a strain field surrounds the solute clusters. As more solute diffuses to the clusters, eventually the matrix can no longer accommodate the matrix mismatch. At this point, a semi-coherent precipitate forms.

Finally, after the semi-coherent precipitate grows to a large enough size, the matrix can no longer support the crystallographic mismatch, and the equilibrium phase forms as incoherent precipitates.

Heating the quenched material in the range of 95-205: C accelerates precipitation in heat treatable alloys. This acceleration is completely not due to changes in reaction rate. As was shown above, structural changes occur that are dependent on time and temperature. In general, the increase in yield strength that occurs during artificial ageing increases faster than the ultimate tensile strength. This means that the alloys lose ductility and toughness. T6 properties are higher than T4 properties, but ductility is reduced. Overageing decreases the tensile strength, and increases the resistance to stress-corrosion-cracking. It also enhances the resistance to fatigue crack growth and imparts dimensional stability to the part.

Precipitation hardening curves have been developed for all the most common alloys. Figure 29.27 shows ageing curves for 2024 and 6061. Both alloys show evidence of reversion of GP zones, as indicated by initial reductions in hardness. This reduction in hardness is caused by the re-solutioning of small GP zones that are below the critical size required for stability. Similar ageing curves have been developed for 7075 and casting alloys.

The ageing curves for different alloys vary; however, generally the higher the ageing temperature, the shorter the time required to attain maximum properties. When high ageing temperatures are used, properties are reached very rapidly. For this reason, ageing temperatures are usually lower to assure that the entire load is brought to the required ageing temperature without risk of reduced properties caused by over-ageing of the fastest heated components.

29.3.7 Appendix: Quenchants

In section 29.3.5 the quenching process step was discussed. In this section, the nature of interactions between the hot part and the quenching medium are discussed and used to explain the choice of quenchants for industrial heat treatment.

For any quenchant, there are typically three phases that occur while the part cools: the vapour phase; nucleate boiling and finally convection. Each of these stages has very specific characteristics and heat transfer mechanisms.

In the vapour phase, a stable gas film of superheated quenchant surrounds the part. The stability of this vapour film depends on several factors, including surface roughness, the boiling temperature of the quenchant and viscosity of the quenchant. Heat transfer is very slow through this film, as heat transfer occurs primarily by radiation and conduction. Because of the relatively low temperatures involved in aluminium alloy solution treatment, radiation heat transfer through the film is negligible. Conduction is also negligible because of the poor conductive heat transfer characteristics of gases.

29-76 H e a t t rea tment

90.0 2024

80.0

70.0 x:"

c 60.0 0 L

�9 50.0 . _

r

40.0 I..-

30.0

20.0 0.1

. f

RT

/ ~ 32oE / ~ , ~ _ ~ - , ~ 0 ~

10 100 1000 10000

Aging time, hours

100000

90.0

6061

80.0

70.0 x:"

e~ 60.0 C O

�9 50.0 . _

t -

O 40.0 }-..

30.0 32~

0OF

RT

20.0 0.1 1 10 100 1000 10000 100000

Aging time, hours

Figure 29.27 Ageing czoa'es /br 2024 amt 6061

As the part cools, the stability of the vapour film also decreases, until the collapse of the vapour film occurs. At this point, nucleate boiling occurs. This transition between stable film boiling, and nucleate boiling is called the 'kedenfrost temperature'.

Nucleate boiling is the fastest regime of cooling during quenching. This is where the vapour stage starts to collapse and all liquid in contact with the component surface erupts into bubbles as boiling occurs. The high heat extraction rates are due to the carrying away of heat from the hot surface and the transferring of it further into the liquid quenchant, which allows cooled liquid to replace it at the surface. In many quenchants, additives have been included to enhance the maximum cooling rates obtained with a given fluid. The boiling stage stops when the temperature of the component's surface reaches a temperature below the boiling point of the liquid. For distortion prone components, high boiling temperature oils or liquid salts could be used, if cooling in these media is fast enough to harden the material, but both of these quenchants see relatively little use in induction hardening.

The final stage of quenching is the convection stage. This occurs when the component has reached a temperature below that of the quenchant's boiling point. Heat is removed by convection and is controlled by the quenchant's specific heat and thermal conductivity, and the temperature differential between the component's temperature and that of the quenchant. The convection stage is usually the slowest of the three quenching stages. Typically, it is this stage where most distortion occurs.

Heat treatment o f aluminium alloys 29-77

900

800

700

o 600 o

'- 500

L

�9 400 E 0

300

200

100

f t Hard water ~, 20~ ~ 4 0 ~ Distilled water .. . . . . 20~ - -o- .40~

' t i ' -

O. & "( " 0 . L~..

'0 -5.. L .

I % t o.

,

k I f ,

. . .

L.

.~ 60~

. . . . . 60~

~ ~ u - O - . @- . . . . - 3 - c ~ - 0 -0- - 0

0 10 20 30 40 50 60

Time, seconds

Figure 29.28 Cooling cula'e oldistilled water showing extended stahle vapour phase

Obtaining the desired properties with low distortion is usually a balancing act. Often, optimal properties are obtained at the expense of high residual stresses or excessive distortion. Conversely, low distortion or residual stresses are usually obtained with a sacrifice in properties. Therefore, the optimum quench rate is one where properties are just met. As a general rule, this usually provides the minimum distortion.

There are several types of quenchants used for heat treating aluminium. The most common are cold water, hot water and polymer quenchants.

29.3.7.1 Cold water quenching

Water is the most common of all quenchant materials. It is easy and inexpensive to obtain, and is readily disposed of, unless severely contaminated. In general, as the temperature of the water is raised, the stability of the vapour phase increases and the onset of nucleate boiling in a stagnant fluid is suppressed. Hence, both the maximum and overall rates of cooling are decreased.

Cold water quenching is the most severe of commonly used quenchants. In an early study using cooling curves, 146 it was shown that quenching into still water caused rapid heat transfer. This study found that heat transfer at the surface of the part was very turbulent at the metal/water interface. The same investigation also showed that there was a marked difference between hard water and distilled water. Distilled water produced an extensive vapour blanket that extended to very low temperatures (Figure 29.28).

29.3.7.2 Hot water quenching

The cooling rate produced by water quenching is independent of material properties, such as thermal conductivity and specific heat. Instead. the cooling rate is primarily dependent on water temperature and agitation. ~47 Water temperature is the largest primary variable controlling the cooling rate. As already discussed, with increasing water temperature, the cooling rate decreases. The maximum cooling rate also decreases, as the water temperature is increased. In addition, the temperature of maximum cooling decreases with increasing water quenchant temperature. The length of time and stability of the vapour barrier increases, with increasing water temperature. This is shown in Table 29.31.

Often, the application of cold water quenchants causes excessive distortion, residual stresses or cracking. Elevated temperatures up to approximately 80=C ( 180 = F) are sometimes used. However, the extended vapour phase that results tends to cause the properties of the workpiece to be compromised, while doing nothing to reduce distortion. In this application, the uniformity of agitation is critical.


Table 29.31 EFFECT OF WATER TEMPERATURE ON COOLING RATES 14s

Temperature corresponding Cooling rate (~ s-! ) at Water Maximum cooling with maximunl cooling workpiece temperature temperature rate rate

(~ (~ s -~ ) (~ 704~ 343~ 232~

40 153 535 60 97 51 50 137 542 32 94 51 60 115 482 20 87 46 70 99 448 17 84 47 80 79 369 15 77 47 90 48 270 12 26 42

Agitation must be able to reach all portions of the part to be effective. Otherwise, vapour can collect and serve as insulation. This will result in soft spots with inadequate properties.

Quenching into water at <50-60~ often produces non-uniform quenching. This non-uniformity manifests itself as spotty hardness, distortion and cracking. This non-uniformity is caused by relatively unstable vapour blanket formation. Because of this difficulty, it was necessary to identify an alternative to water quenching. Polyalkylene glycol quenchants (PAG) were developed to provide a quench rate in between that of water and oil. By control of agitation, temperature and concentration, quench rates similar to water and thick oil can be achieved.

29.3.7.3 Polymer quenching

Polyalkylene glycol polymer quenchants are used in the aerospace industry to control and minimise the distortion occurring during the quenching ofaluminium. Typically these quenchants are governed by AMS 3025, and are either Type I or Type II. Type I quenchants are single polyalkylene glycol polymers, while Type II quenchants are multiple molecular weight polyalkylene glycol polymers. Each offers different benefits. Because of the higher molecular weight of the Type II PAG quenchants, lower concentrations can be used. However, Type II polymers have a lower cloud point temperature, which can cause higher drag-out, if parts are removed from the quenchant before they reach the quenchant temperature, typically 25-40~ (80-100~ Typical concentrations used for Type I polymer quenchants are shown in Table 29.32. A more extensive listing of recommended concentrations is provided in AMS 2770 (Wrought products and forgings) and AMS 2771 (castings).

Polyalkylene glycols exhibit inverse solubility in water. They are completely soluble in water at room temperature, but insoluble at elevated temperatures. The inverse solubility temperature can range from 60~ to 90~ depending on the molecular weight of the polymer, and the structure of the polymer. This phenomenon of inverse solubility modifies the conventional three-stage quenching mechanism 149 and provides great flexibility to control the cooling rate.

When a component is first immersed, the solution in the immediate vicinity of the metal surface is heated to above the inverse solubility temperature. The polymer becomes insoluble and a uniform polymer-rich film encapsulates the surface of the part. The stability and longevity of this polymer film is dependent on the temperature, concentration and amount of agitation present. The stable polymer-rich film eventually collapses uniformly, and cool quenchant comes into contact with the hot metal surface. Nucleate boiling results, with high heat extraction rates.

As the period of active boiling subsides, cooling occurs by conduction and convection into the quenchant. When the surface temperatures fall below the inversion temperatures, the polymer dissolves, forming a homogeneous solution again.

The cooling rate of these polymers can be readily varied to suit the specific application by changing the concentration, quenchant temperature, and the amount of agitation. Typically, for most applications, the agitation is usually fixed, while the concentration is changed.

The concentration of the polymer influences the thickness of the polymer film that is deposited on the surface of the part during quenching. As the concentration increases, the maximum rate of cooling, and the cooling rate in the convection phase decrease.

Agitation has an important effect on the quenching characteristics of the polymer quenchant. It ensures uniform temperature distribution within the quench tank and it also affects the quench rate. As the severity of agitation increases, the lifetime of the polymer-rich phase decreases and eventually disappears, and the maximum rate of cooling increases. Agitation has comparatively little effect on the cooling rate during the convection stage.

Heat treatment o f a luminium alloys

Table 29.32 TYPICAL CONCENTRATION LIMITS FOR QUENCHING IN PAG

2 9 - 7 9

Maximum thickness

Allov Form inches mm Concentration (vol. %1

2024 Sheet 0.040 1.02 0.063 1.60 0.080 2.03

2219 Sheet 0.073 1.85

6061 Sheet, Plate 0.040 1.02 7049 0.190 4.83 7050 7075 0.250 6.35

6061 Forgings 1.0 25 7075 2.0 50

2.5 64

7049 Forgings 3.0 76 7149

7050 Forgings

6061 Extrusions 7049

7050 Extrusions 7075

34 Max. 28 Max. 16 Max. 22 Max.

34 Max. 28 Max.

22 Max.

20-22 13-15 10-12

10-12

3.0 76 20-22

0.250 6.35 28 Max.

0.375 9.52 22 Max.

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2 9 - 8 2 Heat treatment

86.

87.

88.

89.

90.

91.

92.

93.

94.

95.

96.

97. 98.

99.

100.

101.

102.

103.

104.

105.

106. 107.

108.

108a.

109. 110. 111. 112.

113. 114.

115. 116.

117.

118.

119.

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