Steel Metal

Nitriding is a surface-hardening heat treatment that introduces nitrogen into the surface of steel at a temperature range (500 to 550°C, or 930 to 1020°F), while it is in the ferrite condition. Thus, nitriding is similar to carburizing in that surface composition is altered, but different in that nitrogen is added into ferrite instead of austenite. Because nitriding does not involve heating into the austenite phase field and a subsequent quench to form martensite, nitriding can be accomplished with a minimum of distortion and with excellent dimensional control.

The mechanism of nitriding is generally known, but the specific reactions that occur in different steels and with different nitriding media are not always known. Nitrogen has partial solubility in iron. It can form a solid solution with ferrite at nitrogen contents up to about 6%. At about 6% N, a compound called gamma prime (γ’), with a composition of Fe₄N is formed.

At nitrogen contents greater than 8%, the equilibrium reaction product is ε compound, Fe₃N. Nitrided cases are stratified. The outermost surface can be all γ’ and if this is the case, it is referred to as the white layer. Such a surface layer is undesirable: it is very hard profiles but is so brittle that it may spall in use. Usually it is removed; special nitriding processes are used to reduce this layer or make it less brittle. The ε zone of the case is hardened by the formation of the Fe₃N compound, and below this layer there is some solid solution strengthening from the nitrogen in solid solution.

Principal reasons for nitriding are:

• To obtain high surface hardness
• To increase wear resistance
• To improve fatigue life
• To improve corrosion resistance (except for stainless steels)
• To obtain a surface that is resistant to the softening effect of heat at temperatures up to the nitriding temperature

Nitridable Steels

Nitrided steels are generally medium-carbon (quenched and tempered) steels that contain strong nitride-forming elements such as aluminum, chromium, vanadium, and molybdenum.

The most significant hardening is achieved with a class of alloy steels (nitralloy type) that contain about 1% Al. When these steels are nitrided the aluminum forms AlN particles, which strain the ferrite lattice and create strengthening dislocations. Titanium and chromium are also used to enhance case hardness although case depth decreases as alloy content increases.

Of the alloying elements commonly used in commercial steels, aluminum, chromium, vanadium, tungsten and molybdenum are beneficial in nitriding because they form nitrides that are stable at nitriding temperatures. Molybdenum in addition to its contribution as a nitride former also reduces the risk of embrittlement at nitriding temperatures. Other alloying elements such as nickel, copper, silicon and manganese have little, if any, effect on nitriding characteristics.

Although at suitable temperatures all steels are capable of forming iron nitrides in the presence of nascent nitrogen, the nitriding results are more favorable in those steels that contain one or more of the major nitride-forming alloying elements. Because aluminum is the strongest nitride former of the common alloying elements, aluminum containing steels (0.85 to 1.50% Al) yield the best nitriding results in terms of total alloy content.

The following steels can be gas nitrided for specific applications:

1. Aluminum-containing low-alloy steels
2. Medium-carbon, chromium-containing low-alloy steels of the 4100, 4300, 5100, 6100, 8600, 8700 and 9800 series
3. Hot-work die steels containing 5% chromium such as HI1, HI2, and HI3
4. Low-carbon, chromium-containing low-alloy steels of the 3300, 8600, and 9300 series
5. Air-hardening tool steels such as A-2, A-6, D-2, D-3 and S-7
6. High-speed tool steels such as M-2 and M-4
7. Nitronic stainless steels such as 30, 40, 50, and 60
8. Ferritic and martensitic stainless steels of the 400 and 500 series
9. Austenitic stainless steels of the 200 and 300 series
10. Precipitation-hardening stainless steels such as 13-8 PH, 15-5 PH, 17-4 PH, 17-7 PH, A-286, AM350 and AM355.

Nitriding processes

Process methods for nitriding include:

• gas (box furnace or fluidized bed),
• liquid (salt bath),
• plasma (ion) nitriding.

The advantages and disadvantages of these techniques are similar to those of carburizing. However, times for gas nitriding can be quire long, that is, from 10 to 130 h depending on the application, and the case depths are relatively shallow, usually less than 0.5 mm. Plasma nitriding allows faster nitriding times, and the quickly attained surface saturation of the plasma process results in faster diffusion. Plasma nitriding can also clean the surface by sputtering.

Gas Nitriding

Gas nitriding is a case-hardening process whereby nitrogen is introduced into the surface of a solid ferrous alloy by holding the metal at a suitable temperature in contact with a nitrogenous gas, usually ammonia. Quenching is not required for the production of a hard case. The nitriding temperature for all steels is between 495 and 565°C.

Because of the absence of a quenching requirement with attendant volume changes, and the comparatively low temperatures employed in this process, nitriding of steels produces less distortion and deformation than either carburizing or conventional hardening. Some growth occurs as a result of nitriding but volumetric changes are relatively small.

Prior Heat Treatment. All hardenable steels must be hardened and tempered before being nitrided. The tempering temperature must be high enough to guarantee structural stability at the nitriding temperature: the minimum tempering temperature is usually at least 30°C (50°F) higher than the maximum temperature to be used in nitriding.

Single-Stage and Double-Stage Nitriding. Either a single- or a double-stage process may be employed when nitriding with anhydrous ammonia. In the single-stage process, a temperature in the range of about 495 to 525°C is used and the dissociation rate ranges from 15 to 30%. This process produces a brittle nitrogen-rich layer known as the white nitride layer at the surface of the nitrided case.

The double-stage process, known also as the Floe process, has the advantage of reducing the thickness of the white nitrided layer.

The first stage of the double-stage process is, except for time, a duplication of the single-stage process. The second stage may proceed at the nitriding temperature employed for the first stage or the temperature may be increased to from 550 to 565°C; however, at either temperature, the rate of dissociation in the second stage is increased to 65 to 80% (preferably 75 to 80%). Generally, an external ammonia dissociator is necessary for obtaining the required higher second-stage dissociation.

The principal purpose of double-stage nitriding is to reduce the depth of the white layer produced on the surface of the case. Except for a reduction in the amount of ammonia consumed per hour, there is no advantage in using the double-stage process unless the amount of white layer produced in single-stage nitriding cannot be tolerated on the finished part or unless the amount of finishing required after nitriding is substantially reduced.

To summarize, the use of a higher temperature during the second stage:

• Lowers the case hardness
• Increases the case depth
• May lower the core hardness depending on the prior tempering temperature and the total nitriding cycle time
• May lower the apparent effective case depth because of the loss of core hardness depending on how effective case depth is defined.

Operating Procedures. After hardening and tempering and before nitriding, parts should be thoroughly cleaned. Most parts can be successfully nitrided immediately after vapor degreasing.

Bright Nitriding

Bright nitriding is a modified form of gas nitriding employing ammonia and hydrogen gases. Atmosphere gas is continually withdrawn from the nitriding furnace and passed through a temperature-controlled scrubber containing a water solution of sodium hydroxide (NaOH). Trace amounts of hydrogen cyanide (HCN) formed in the nitriding furnaces are removed in the scrubber thus improving the rate of nitriding.

The scrubber also establishes a predetermined moisture content in the nitriding atmosphere reducing the rate of cyanide formation and inhibiting the cracking of ammonia to molecular nitrogen and hydrogen. By this technique control over the nitrogen activity of the furnace atmosphere is enhanced and nitrided parts can be produced with little or no white layer at the surface. If present, the white layer will be composed of only the more ductile Fe₄N (gamma prime) phase.

Pack Nitriding

Pack nitriding is a process analogous to pack carburizing. It employs certain nitrogen-bearing organic compounds as a source of nitrogen. Upon heating, the compounds used in the process form reaction products that are relatively stable at temperatures up to 570°C.

Slow decomposition of the reaction products at the nitriding temperature provides a source of nitrogen. Nitriding times of 2 to 16 h can be employed. Pans are packed in glass ceramic or aluminum containers with the nitriding compound, which is often dispersed in an inert packing media.

Ion (or Plasma) Nitriding

Since the mid-1960s, nitriding equipment utilizing the glow-discharge phenomenon has been commercially available. Initially termed glow-discharge nitriding, the process is now generally known as ion, or plasma, nitriding. The term plasma nitriding is gaining acceptance.

Ion nitriding is an extension of conventional nitriding processes using plasma-discharge physics. In vacuum, high-voltage electrical energy is used to form a plasma, through which nitrogen ions are accelerated to impinge on the workpiece. This ion bombardment heats the workpiece, cleans the surface, and provides active nitrogen.

Metallurgically versatile, the process provides excellent dimensional control and retention of surface finish. Ion nitriding can be conducted at temperatures lower than those conventionally employed. Control of white-layer composition and thickness enhances fatigue properties. The span of ion-nitriding applications includes conventional ammonia- gas nitriding, short-cycle nitriding in salt bath or gas, and the nitriding of stainless steels.

Ion nitriding lends itself to total process automation, ensuring repetitive metallurgical results. The absence of pollution and insignificant gas consumption are important economic and public policy factors. Moreover, selective nitriding accomplished by simple masking techniques may yield significant economies.

Comparison of Ion Nitriding and Ammonia-Gas Nitriding

Ammonia-gas nitriding produces a compound zone that is a mixture of both epsilon and gamma-prime structures. High internal stresses result from differences in volume growth associated with the formation of each phase. The interfaces between the two crystal structures are weak. Thicker compound zones, formed by ammonia-gas nitriding, limit accommodation of the internal stresses resulting from the mixed structure.

Under cyclic loading, cracks in the compound zone can serve as initiation points for the propagation of fatigue cracks. The single-phase gamma-prime compound zone, which is thin and more ductile, exhibits superior fatigue properties. Reducing the thickness of the ion-nitrided compound zone further improves fatigue performance. Maximization occurs at the limiting condition, where compound zone depth equals zero.

Case Hardness. The bulk of the thickness of the nitride case is the diffusion zone where fine iron/alloy nitride precipitates impart increased hardness and strength. Compressive stresses are also developed, as in other nitriding processes. Hardness profiles resulting from ion nitriding are similar to ammonia-gas nitriding but near-surface hardness may be greater with ion nitriding, a result of lower processing temperature.

Advantages and Disadvantages of Ion Nitriding. Ion nitriding achieves repetitive metallurgical results and complete control of the nitrided layers. This control results in superior fatigue performance, wear resistance, and hard layer ductility. Moreover, the process ensures high dimensional stability, eliminates secondary operations, offers low operating-temperature capability and produces parts that retain surface finish.

Among operating benefits are:

• Total absence of pollution
• Efficient use of gas and electrical energy
• Total process automation
• Selective nitriding by simple masking techniques
• Process span that encompasses all sub-critical nitriding
• Reduced nitriding time

The limitations of ion nitriding include high capital cost, need for precision fixturing with electrical connections, long processing times compared to other short-cycle nitrocarburizing processes, and lack of feasibility of liquid quenching for carbon steels.

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Carburizing is the addition of carbon to the surface of low-carbon steels at temperatures generally between 850 and 950°C (1560 and 1740°F), at which austenite, with its high solubility for carbon, is the stable crystal structure. Hardening is accomplished when the high-carbon surface layer is quenched to form martensite so that a high-carbon martensitic case with good wear and fatigue resistance is superimposed on a tough, low-carbon steel core.

Case hardness of carburized steels is primarily a function of carbon content. When the carbon content of the steel exceeds about 0.50% additional carbon has no effect on hardness but does enhance hardenability. Carbon in excess of 0.50% may not be dissolved, which would thus require temperatures high enough to ensure carbon-austenite solid solution.

Case depth of carburized steel is a function of carburizing time and the available carbon potential at the surface. When prolonged carburizing times are used for deep case depths, a high carbon potential produces a high surface-carbon content, which may thus result in excessive retained austenite or free carbides. These two microstructural elements both have adverse effects on the distribution of residual stress in the case-hardened part. Consequently, a high carbon potential may be suitable for short carburizing times but not for prolonged carburizing.

Carburizing steels for case hardening usually have base-carbon contents of about 0.2%, with the carbon content of the carburized layer generally being controlled at between 0.8 and 1% C. However, surface carbon is often limited to 0.9% because too high a carbon content can result in retained austenite and brittle martensite.

Most steels that are carburized are killed steels (deoxidized by the addition of aluminum), which maintain fine grain sizes to temperatures of about 1040°C. Steels made to coarse grain practices can be carburized if a double quench provides grain refinement. Double quenching usually consists of a direct quench and then a requench from a lower temperature.

Many alloy steels for case hardening are now specified on the basis of core hardenability. Although the same considerations generally apply to the selection of uncarburized grades, there are some peculiarities in carburizing applications.

First, in a case-hardened steel, the hardenability of both case and core must be considered. Because of the difference in carbon content, case and core have quite different hardenabilities, and this difference is much greater for some steels than for others.

Moreover, the two regions have different in-service functions to perform. Until the introduction of lean alloy steels such as the 86xx series, with and without boron, there was little need to be concerned about case hardenability because the alloy content combined with the high carbon content always provided adequate hardenability. This is still generally true when the steels are direct quenched from carburizing, so that the carbon and alloying elements are in solution in the case austenite. In parts that are reheated for hardening and in heavy-sectioned parts, however, both case and core hardenability requirements should be carefully evaluated.

The relationship between the thermal gradient and the carbon gradient during quenching of a carburized part can make a measurable difference in the case depth as measured by hardness. That is, an increase in base hardenability can produce a higher proportion of martensite for a given carbon level, yielding an increased measured case depth. Therefore, a shallower carbon profile and shorter carburizing time could be used to attain the desired result in a properly chosen steel.

Core Hardness. A common mistake is to specify too narrow a range of core hardness. When the final quench is from a temperature high enough to allow the development of full core hardness, the hardness variation at any location will be that of the hardenability band of the steel at the corresponding position on the end-quenched hardenability specimen.

In standard steels purchased to chemical composition requirements rather than to hardenability, the range can be 20 or more HRC points; for example, 8620 may vary from 20 to 45 HRC at the 4/16 in.(6.35mm) position. The 25-point range emphasizes the advantage of purchasing to hardenability specifications to avoid the intolerable variation possible within the ranges for standard chemistry steels. Another way to control core hardness within narrow limits without resorting to the use of high-alloy steels is to use a final quench from a lower temperature so that full hardness in the case will be developed without the disadvantage of excessive core hardness.

Gears are almost always oil quenched because distortion must be held to the lowest possible level. Therefore, alloy steels are usually selected, with much debate about which particular alloy. The lower-alloy steels such as 4023, 5120, 4118, 8620, and 4620, with a carbon range between 0.15 and 0.25%, are widely used and generally satisfactory. Usually, the first choice is one of the last two steels mentioned, either of which should be safe for all ordinary applications. The final choice, based on service experience or dynamometer testing, should be the least expensive steel that will do the job. For heavy-duty applications, higher-alloy grades such as 4320, 4817, and 9310 are justifiable if based on actual performance tests. The life testing of gears in the same mountings used in service to prove both the design and the steel selection is particularly important.

In other applications, when distortion is not a major factor, the carbon steels described above, water quenched, can be used up to a 50 mm (2 in.) diameter. In larger sizes, low-alloy steels, water quenched, such as 5120, 4023, and 6120 can be used, but possible distortion and quench cracking must be avoided.

Carburizing Methods. While the basic principle of carburizing has remained unchanged since carburizing was first employed, the method used to introduce the carbon into the steel has been a matter of continuous evolution.

In its earliest application, parts were simply placed in a suitable container and covered with a thick layer of carbon powder (pack carburizing). Although effective in introducing carbon, this method was exceedingly slow, and as the demand for greater production grew, a new process using a gaseous atmosphere was developed.

In gas carburizing, the parts are surrounded by a carbon-bearing atmosphere that can be continuously replenished so that a high carbon potential can be maintained. While the rate of carburizing is substantially increased in the gaseous atmosphere, the method requires the use of a multicomponent atmosphere whose composition must be very closely controlled to avoid deleterious side effects, for example, surface and grain-boundary oxides. In addition, a separate piece of equipment is required to generate the atmosphere and control its composition. Despite this increased complexity, gas carburizing has become the most effective and widely used method for carburizing steel parts in large quantities.

In efforts required to simplify the atmosphere, carburizing in an oxygen-free environment at very low pressure (vacuum carburizing) has been explored and developed into a viable and important alternative. Although the furnace enclosure in some respects becomes more complex, the atmosphere is greatly simplified. A single-component atmosphere consisting solely of a simple gaseous hydrocarbon, for example methane, may be used. Furthermore, because the parts are heated in an oxygen-free environment, the carburizing temperature may be increased substantially without the risk of surface or grain-boundary oxidation. The higher temperature permitted increases not only the solid solubility of carbon in the austenite but also its rate of diffusion, so that the time required to achieve the case depth desired is reduced.

Although vacuum carburizing overcomes some of the complexities of gas carbunzing, it introduces a serious new problem that must be addressed. Because vacuum carburizing is conducted at very low pressures, and the rate of flow of the carburizing gas into the furnace is very low, the carbon potential of the gas in deep recesses and blind holes is quickly depleted. Unless this gas is replenished, a great nonuniformity in case depth over the surface of the part is likely to occur. If, in an effort to overcome this problem, the gas pressure is increased significantly, another problem arises, that of free-carbon formation, or sooting.

Thus, in order to obtain cases of reasonably uniform depth over a part of complex shape, the gas pressure must be increased periodically to replenish the depleted atmosphere in recesses and then reduced again to the operating pressure. Clearly, a delicate balance exists in vacuum carburizing: The process conditions must be adjusted to obtain the best compromise between case uniformity, risk of sooting, and carburizing rate.

A method that overcomes both of these major problems, yet retains the desirable features of a simple atmosphere and permissible operating temperature is plasma or ion carburizing.

To summarize, carburizing methods include:

• Gas carburizing
• Vacuum carburizing
• Plasma carburizing
• Salt bath carburizing
• Pack carburizing

These methods introduce carbon by the use of gas (atmospheric-gas, plasma, and vacuum carburizing), liquids (salt bath carburizing), or solid compounds (pack carburizing). All of these methods have limitations and advantages, but gas carburizing is used most often for large-scale production because it can be accurately controlled and involves a minimum of special handling.

Vacuum carbunzing and plasma carburizing have found applications because of the absence of oxygen in the furnace atmosphere. Salt bath and pack carburizing arc still done occasionally, but have little commercial importance today.

Process characteristics of the above-mentioned carburizing methods fall into two general groups:

• Conventional methods, which introduce carbon by gas atmospheres, salt baths or charcoal packs
• Plasma methods, which impinge positive carbon ions on the surface of a steel part (the cathode)

The main difference between conventional and plasma methods is the reduced carburizing times achieved in plasma-assisted methods. The quickly attained surface saturation also results in faster diffusion kinetics. Furthermore, plasma carburizing produces very uniform case depths, even in parts with irregular surfaces.

With the conventional methods, carburization always takes place by means of a gaseous phase of carbon monoxide; however, each method also involves different reaction and surface kinetics, producing different case-hardening results.

In general, with conventional methods, carbon monoxide breaks down at the steel surface:

2CO ↔ CO₂ + C

The liberated carbon is readily dissolved by the austenite phase and diffuses into the body of the steel. For some process methods (gas and pack carburizing), the carbon dioxide produced may react with the carbon atmosphere or pack charcoal to produce new carbon monoxide by the reverse reaction.

Carburizing is most frequently performed between 850 and 950°C (1550 and 1750°F), but sometimes higher temperatures are used to reduce cycle times and/or produce deeper depths of the high-carbon surface layer.

A comprehensive model of gas carburization must include algorithms that describe:

• Carbon diffusion
• Kinetics of the surface reaction
• Kinetics of the reaction between endogas and enriching gas
• Purging (for batch processes)
• The atmosphere control system.

Surface hardening a process which includes a wide variety of techniques is used to improve the wear resistance of parts without affecting the softer, tough interior of the part. This combination of hard surface and resistance and breakage upon impact is useful in parts such as a cam or ring gear that must have a very hard surface to resist wear, along with a tough interior to resist the impact that occurs during operation. Further, the surface hardening of steels has an advantage over through hardening because less expensive low-carbon and medium-carbon steels can be surface hardened without the problems of distortion and cracking associated with the through hardening of thick sections.

There are two distinctly different approaches to the various methods for surface hardening (Table 1): methods that involve an intentional buildup or addition of a new layer and methods that involve surface and subsurface modification without any intentional buildup or increase in part dimensions.

Table 1. Engineering methods for surface hardening of steels.

Layer additions	Substrate treatment
Hardfacing Fusion harcifacing Thermal spray Coatings Electrochemical plating Chemical vapor deposition (electroless plating) Thin films (physical vapor deposition, puttering, ion plating) Ion mixing	Diffusion methods Carburizing Nitriding Carbonitriding Nitrocarburizing Boriding Titanium-carbon diffusion Toyota diffusion process Selective hardening methods Flame hardening Induction hardening Laser hardening Electron beam hardening Ion implantation Selective carburizing and nitriding Use of arc lamps

The first group of surface hardening methods includes the use of thin films, coatings, or weld overlays (hard-facings). Films, coatings, and overlays generally become less cost effective as production quantities increase, especially when the entire surface of work pieces must be hardened.

The fatigue performance of films, coatings, and overlays may also be a limiting factor, depending on the bond strength between the substrate and the added layer. Fusion-welded overlays have strong bonds, but the primary surface-hardened steels used in wear applications with fatigue loads include heavy case-hardened steels and flame or induction-hardened steels. Nonetheless, coatings and overlays can be effective in some applications. For tool steels, for example, TiN and Al2O3 coatings are effective not only because of their hardness but also because their chemical inertness reduces wear and the welding of chips to the tool. Overlays can be effective when the selective hardening of large areas is required.

The second group of methods on surface hardening is further divided into diffusion methods and selective hardening methods. Diffusion methods modify the chemical composition of the surface with hardening species such as carbon, nitrogen, or boron. Diffusion methods allow effective hardening of the entire surface of a part and are generally used when a large number of parts are to be surface hardened. In contrast, selective surface hardening methods allow localized hardening. Selective hardening generally involves transformation hardening (from heating and quenching), but some selective hardening methods (selective nitriding, ion implantation and ion beam mixing) are based solely on compositional modification.

As previously mentioned, surface hardening by diffusion involves the chemical modification of a surface. The basic process used is thermo-chemical because some heat is needed to enhance the diffusion of hardening species into the surface and subsurface regions of part.

The depth of diffusion exhibits time-temperature dependence such that:

Case depth ≈ K √Time

where the diffusivity constant, K, depends on temperature, the chemical composition of the steel, and the concentration gradient of a given hardening species. In terms of temperature, the diffusivity constant increases exponentially as a function of absolute temperature. Concentration gradients depend on the surface kinetics and reactions of a particular process.

Methods of hardening by diffusion include several variations of hardening species (such as carbon, nitrogen, or boron) and of the process method used to handle and transport the hardening species to the surface of the part. Process methods for exposure involve the handling of hardening species in forms such as gas, liquid, or ions. These process variations naturally produce differences in typical case depth and hardness (Table 2). Factors influencing the suitability of a particular diffusion method include the type of steel (Table 3).

It is also important to distinguish between total case depth and effective case depth. The effective case depth is typically about two-thirds to three-fourths the total case depth. The required effective depth must be specified so that the heat treatment can process the parts for the correct time at the proper temperature.

Table 2: Typical characteristics of diffusion treatments

Process	Nature of case	Process temperature (°C)	Typical case depth	Case hardness (HRC)	Typical base metals
Carburizing Pack	Diffused carbon	815-1090	125μm-1.5mm	50-63*	Low-carbon steels, low-carbon alloy steels
Gas	Diffused carbon	815-980	75 μm-1.5mm	50-63*	Low-carbon steels, low-carbon alloy steels
Liquid	Diffused carbon and possibly nitrogen	815-980	50 μm-1.5mm	50-65*	Low-carbon steels, low-carbon alloy steels
Vacuum	Diffused carbon	815-1090	75 μm-1.5mm	50-63*	Low-carbon steels, low-carbon alloy steels
Nitriding Gas	Diffused nitrogen, nitrogen compounds	480-590	12μm-0.75mm	50-70	Alloy steels, nitriding steels, stainless steels
Salt	Diffused nitrogen, nitrogen compounds	510-565	2.5μm-0.75mm	50-70	Most ferrous metals. Including cast irons
Ion	Diffused nitrogen. nitrogen compounds	340-565	75μm-0.75mm	50-70	Alloy steels, nitriding steels, stainless steels
Carbonitriding Gas	Diffused carbon and nitrogen	760-870	75μm-0.75mm	50-65*	Low-carbon steels, low-carbon alloy steels, stainless steels
Liquid (cyaniding)	Diffused carbon and nitrogen	760-870	2.5-125μm	50-65*	Low-carbon steels
Ferritic nitrocarburizing	Diffused carbon and nitrogen	565-675	2.5-25μm	40-60*	Low-carbon steels
Other Aluminizing (pack)	Diffused aluminum	870-980	25μm-1mm	< 20	Low-carbon steels
Siliconizing by chemical vapor deposition	Diffused silicon	925-1040	25μm-1mm	30-50	Low-carbon steels
Chromizing by chemical vapor deposition	Diffused chromium	980-1090	25-50μm	Low-carbon steel < 30; High-carbon 50-60	High- and low carbon steels
Titanium Carbide	Diffused carbon and titanium, TiC compound	900-1010	2,5-12.5μm	> 70*	Alloy steels, tool steels
Boriding	Diffused boron. boron compounds	400-1150	12,5-50μm	40- > 70	Alloy steels, tool steels,Cobalt and nickel alloys

*  Requires quench from austenitizing temperature.
Table 3. Types of steels used for various diffusion processes

Diffusion substrates
Low-carbon steels	Alloy steels	Tool steels	Stainless steels
Carburizing Cyaniding Ferritic nitrocarburizing Carbonitriding	Nitriding Ion nitriding	Titanium carbide Boriding Salt nitriding Ion nitriding Gas nitriding	Gas nitriding Titanium carbide Ion nitriding Ferritic nitrocarburizing

Thermomechanical treatment involves the simultaneous application of heat and a deformation process to an alloy, in order to change its shape and refine the microstructure. Thus, hot-rolling of metals, a well-established industrial process, is a thermomechanical treatment which plays an important part in the processing of many steels from low carbon, mild steels to highly alloyed stainless steels. The traditional fabrication route involves the casting of ingots varying in size from 1 to 50 tones, which are soaked at very high temperatures (1200-1300°C), then progressively hot rolled to billets, bars and sheet. This leads to the breaking down of the original coarse cast structure by repeated recrystallization of the steel while in the austenitic condition, and by the gradual reduction of inhomogeneities of composition caused by segregation during casting. Also, the inevitable non-metallic inclusions, i.e. oxides, silicates, sulphides, are broken up, some deformed, and distributed throughout the steel in a more uniform manner.

Low temperature thermomechanical treatment -LTMT (Ausforming)

The process known as ausforming or low temperature thermomechanical treatment (LTMT), involves the deformation of austenite in the metastable bay between the ferrite and bainite curves of the TTT diagram. The treatment is shown schematically in Fig. 1a. Steel, with a sufficiently developed metastable austenite bay is quenched from the austenitizing temperature to this region, where substantial deformation is carried out, without allowing transformation to take place. The deformed steel is then transformed to martensite during quenching to room temperature, and the appropriate balance of mechanical properties achieved by subsequent tempering. This ausforming treatment can be contrasted with a high temperature thermomechanical treatment (HTMT), where the deformation is carried out in the stable austenite region (Fig. 1b), usually above Ac3 prior to quenching to form martensite. In a third process, isoforming (Fig. 1c), the steel is deformed in the metastable austenite region, but the deformation is continued until the transformation is complete at the intermediate temperature. The steel can then be slowly cooled to room temperature.

The ausforming process needs careful control to be successful and usually involves very substantial deformation. However, the attraction is that with appropriate steels dramatic increases in strength are achieved without adverse effect on ductility and toughness. Typically, a 4,7% Cr, 1.5%Mo, 0.4%V, 0.34%C steel has a tensile strength of about 2000 MPa after conventional quenching and tempering, whereas after ausforming the strength can be over 3000 MPa.

Steels, in which austenite transforms rapidly at subcritical temperatures, are not suitable for ausforming. It is necessary to add alloying elements which develop a deep metastable austenite bay by displacing the TTT curve to longer transformation times. The most useful elements in this respect are chromium, molybdenum, nickel and manganese, and allowance must be made for the fact that deformation of the austenite accelerates the transformation. Consequently, it is necessary to have sufficient alloying element present to slow down the reaction and avoid the formation of ferrite during cooling to the deformation temperature.

Figure 1. Schematic diagrams of thermochemical treatments:
a) ausforming-low temperature mechanical treatment;
b) high temperature mechanical treatment;
c) isoforming transformation.

The austenite grain size should be as tine as possible, not only to increase the dislocation density during deformation but also to minimize the martensite plate size on quenching from the metastable austenite bay.

Cooling from the austenitizing temperature to the metastable bay must be sufficiently rapid to avoid the formation of ferrite and, after deformation, the cooling should be fast enough to prevent the formation of bainite. The strength achieved as a result of ausforming increases as the deformation temperature is decreased, presumably because of the greater strain hardening induced in the austenite. In any case, the temperature chosen should be low enough to avoid recovery and recrystallization, but high enough to prevent bainite from forming during the deformation. Premature austenite decomposition has been found to be detrimental to mechanical properties.

The amount of deformation is a most important variable. There is a roughly linear relationship between the degree of working and the strength finally achieved, with increases between 4 and 8 MPa per percent deformation. One of the most significant trends is that for many steels the ductility actually increases with increasing deformation, although this only becomes significant at deformations above 30% reduction in thickness.

As might be expected, steels subjected to heavy deformation during ausforming exhibit very high dislocation densities (up to 10¹³cm^-2) formed partly during deformation and partly during the shear transformation to martensite. The deformation is usually carried out in the temperature range (500-600°C) in which alloy carbides would be expected to precipitate, so it is not surprising that fine alloy carbide dispersions have been detected by dark field electron microscopy.

On transforming the warm worked austenite to martensite, it is likely that at least part of the dislocation substructure, together with the fine carbide dispersion, is inherited by the martensite. The martensite plate size has been shown to be very substantially smaller than in similar steels given a straight quench from the austenitizing temperature.

Several factors must contribute to strength because anyone mechanism cannot fully account for the high degree of strengthening observed. However, it seems likely that the major contributions are from the very high dislocation density and the fine dispersion of alloy carbides associated with the dislocations. It should also be added that the fine precipitate particles can act as dislocation multiplication centers during plastic deformation. The martensitic transformation is an essential part of the strengthening process, as it substantially increases the dislocation density and divides each deformed austenite grain into a large number of martensitic plates, which are much smaller than those in conventional heat treatments. It is also likely that these small plates have inherited fine dislocation substructures from the deformed metastable austenite.

Isoforming

The process of isoforming involves deformation of metastable austenite, but the deformation is continued until the transformation of austenite is complete at the deformation temperature (Fig. 1c). This is because the lamellar morphology of pearlite leads to low toughness in ferrite/pearlite steels, the ductile/brittle transition temperature increasing with larger volume fraction of pearlite. However, by applying deformation during the phase transformation, instead of a ferrite/pearlite aggregate, the structure produced consists of fine ferrite subgrains (≈0.5μm diameter) with spheroidized cementite particles (≈25nm diameter) mainly located at subgrain triple points.

As in the case of steels for ausforming, the chosen steel must have a suitable TTT diagram. First, it is necessary to be able to deform the austenite prior to transformation, then the transformation must be complete before deformation has ceased. Only modest increases in strength are achieved. However, there can be a very substantial improvement in toughness due to the refinement of the ferrite grain size and the replacement of lamellar cementite by spheroidized particles. However, for significant gains in toughness, deformations in excess of 70% reduction in area are needed. Finally, care must be taken to restrict deformation to temperatures at which the ferrite and pearlite reactions take place as similar deformation in the bainitic region leads to marked reductions in toughness.

High temperature thermomechanical treatments (HTMT)

In high temperature thermomechanical treatments the deformation is carried out in the stable austenite range just above Ac₃ (Fig. 1b), and so can be performed in steels, which do not possess a suitable metastable austenite bay. The steel is then quenched to the martensitic state and tempered at an appropriate temperature. The strengthening achieved arises from austenite grain size refinement, typically from 10-60 μm to 3 μm, but optimum properties are often obtained if recryslallization of the austenite is avoided. As in ausforming strong carbide forming elements are beneficial, which suggests that alloy carbide precipitation occurs in the austenite during deformation. A particular advantage of this process is that optimum properties can be achieved at modest deformations (≈40%) so that deformation can be carried out more readily on existing equipment. The HTMT process does not yield as high strengths as in ausforming but the ductility and fatigue properties are usually superior.

Clearly, HTMT is a variant of controlled rolling. However, it is normally applied to steels with higher alloying contents which can then be transformed to martensite and tempered.

Industrial steels subjected to thermomechanical treatments

Ausforming has provided some of the strongest, toughest steels so far produced, with the added advantage of very good fatigue resistance. However, they usually have high concentrations of expensive alloying elements and must be subjected to large deformations, which impose heavy workloads on rolling mills. Nevertheless, these steels are particularly useful where a high strength to weight ratio is required and where cost is a secondary factor. Typical applications have included parts for undercarriages of aircraft, special springs and bolts.

The 12%Cr transformable steels respond readily to ausforming to the extent that tensile strengths of over 3000MPa can be obtained in appropriate compositions. 0.4C-6Mn-3Cr-1.5Si steel has been ausformed to a tensile strength of 3400 MPa, with an improvement in ductility over the conventional heat treatment. Similar high strength levels with good ductility have been reported for 0.4C-5Cr-1.3Mo-1.0Si-0.5V steel. All of these steels are sufficiently highly alloyed to allow adequate time for substantial deformation in the austenite bay of the TTT curve prior to transformation.

Although pure iron is a weak material, steels cover a wide range of the strength spectrum from low yield stress levels (around 200 MPa) to very high levels (approaching 2000 MPa). These mechanical properties are usually achieved by the combined use of several strengthening mechanisms, and in such circumstances it is often difficult to quantify the different contributions to the strength. These results should then be helpful in examining the behavior of more complex steels.

Like other metals, iron can be strengthened by several basic mechanisms, the most important of which are:

• Work hardening
• Solid solution strengthening by interstitial atoms
• Solid solution strengthening by substitutional atoms
• Refinement of grain size
• Dispersion strengthening, including lamellar and random dispersed structures.

The most distinctive aspect of strengthening of iron is the role of the interstitial solutes carbon and nitrogen. These elements also play a vital part in interacting with dislocations, and in combining preferentially with some of the metallic alloying elements used in steels.

Work hardening

Work hardening is an important strengthening process in steel, particularly in obtaining high strength levels in rod and wire, both in plain carbon and alloy steels. For example, the tensile strength of a 0.05% C steel subjected to 95% reduction in area by wire drawing, is raised by no less than 550 MPa while higher carbon steels are strengthened by up to twice this amount. Indeed, without the addition of special alloying elements, plain carbon steels can be raised to strength levels above 1500 MPa simply by the phenomenon of work hardening.

Basic work on the deformation of iron has largely concentrated on the other end of the strength spectrum, namely pure single crystals and polycrystals subjected to small controlled deformations. The diversity of slip planes leads to rather irregular wavy slip bands in deformed crystals, as the dislocations can readily move from one type of plane to another by cross slip, provided they share a common slip direction.

The yield stress of iron single crystals are very sensitive to both temperature and strain rate and a similar dependence has been found for less pure polycrystalline iron. Therefore, the temperature sensitivity cannot be attributed to interstitial impurities. It is explained by the effect of temperature on the stress needed to move free dislocations in the crystal, the Peierls-Nabarro stress.

Solid solution strengthening by interstitials

The formation of interstitial atmospheres at dislocations requires diffusion of the solute. As both carbon and nitrogen diffuse much more rapidly in iron than substitutional solutes, it is not surprising that strain ageing can take place readily in the range from 20°C to 150°C. Consequently the atmosphere condenses to form rows of interstitial atoms along the cores of the dislocations. These arise because the temperature is high enough to allow interstitial atoms to diffuse during deformation, and to form atmospheres around dislocations generated throughout the stress-strain curve. Steels tested under these conditions also show low ductility`s, due partly to the high dislocation density and partly to the nucleation of carbide particles on the dislocations where the carbon concentration is high. The phenomenon is often referred to as blue brittleness, blue being the interference color of the steel surface when oxidized in this temperature range.

The break away of dislocations from their carbon atmospheres as a cause of the sharp yield point became a controversial aspect of the theory because it was found that the provision of free dislocations, for example, by scratching the surface of a specimen, did not eliminate the sharp yield point. An alternative theory was developed which assumed that, once condensed carbon atmospheres are formed in iron, the dislocations remain locked, and the yield phenomena arise from the generation and movement of newly formed dislocations.

To summarize, the occurrence of a sharp yield point depends on the occurrence of a sudden increase in the number of mobile dislocations. However, the precise mechanism by which this takes place will depend on the effectiveness of the locking of the pre-existing dislocations. If the pinning is weak, then the yield point can arise as a result of unpinning. However, if the dislocations are strongly locked, either by interstitial atmospheres or precipitates, the yield point will result from the rapid generation of new dislocations.

Under conditions of dynamic strain ageing, where atmospheres of carbon atoms form continuously on newly-generated dislocations, it would be expected that a higher density of dislocations would be needed to complete the deformation, if it is assumed that most dislocations which attract carbon atmospheres are permanently locked in position.

Strengthening at high interstitial concentrations

Austenite can take into solid solution up to 10% carbon, which can be retained in solid solution by rapid quenching. However, in these circumstances the phase transformation takes place, not to ferrite but to a tetragonal structure referred to as martensite. This phase forms as a result of diffusion less shear transformation leading to characteristic laths or plates.

If the quench is sufficiently rapid, the martensite is essentially a supersaturated solid solution of carbon in a tetragonal iron matrix, and as the carbon concentration can be greatly in excess of the equilibrium concentration in ferrite, the strength is raised very substantially. High carbon martensites are normally very hard but brittle, the yield strength reaching as much as 1500 MPa; much of this increase can be directly attributed to increased interstitial solid solution hardening, but there is also a contribution from the high dislocation density, which is characteristic of martensitic transformations in iron-carbon alloys.

Substitutional solid solution strengthening of iron

Many metallic elements form solid solutions in γ- and α-iron. These are invariably substitutional solid solutions, but for a constant atomic concentration of alloying elements there are large variations in strength. Using single crystal data for several metals, Fig. 1 shows that an element such as vanadium has a weak strengthening effect on α-iron at low concentrations (< 2%), while silicon and molybdenum are much more effective strengthened. Other data indicates that phosphorus; manganese, nickel and copper are also effective strengtheners. However, it should be noted that the relative strengthening might alter with the temperature of testing, and with the concentrations of interstitial solutes present in the steels.

Figure 1. Solid solution strengthening of iron crystals by substitutional solutes. Ratio of the critical resolved shear stress τ₀ to shear modulus μ as a function of atomic concentration.

The strengthening achieved by substitutional solute atoms is, in general, greater the larger the difference in atomic size of the solute from that of iron, applying the Hume-Rothery size effect. However, from the work of Fleischer and Takeuchi it is apparent that differences in the elastic behavior of solute and solvent atoms are also important in determining the overall strengthening achieved.

In practical terms, the contribution to strength from solid solution effects is superimposed on hardening from other sources, e.g. grain size and dispersions. Also it is a strengthening increment, like that due to grain size, which need not adversely affect ductility. In industrial steels, solid solution strengthening is a far from negligible factor in the overall strength, where it is achieved by a number of familiar alloying elements, e.g. manganese, silicon, nickel, molybdenum, several of which are frequently present in a particular steel and are additive in their effect. These alloying elements arc usually added for other reasons, e.g. Si to achieve deoxidation, Mn to combine with sulphur or Mo to promote hardenability. Therefore, the solid solution hardening contribution can be viewed as a useful bonus.

Grain size

The refinement of the grain size of ferrite provides one of the most important strengthening routes in the heal treatment of steels. The grain size effect on the yield stress can therefore be explained by assuming that a dislocation source operates within a crystal causing dislocations to move and eventually to pile up at the grain boundary. The pile-up causes a stress to be generated in the adjacent grain, which, when it reaches a critical value, operates a new source in that grain.

In this way, the yielding process is propagated from grain to grain. The grain size determines the distance dislocations have to move to form grain boundary pile-ups, and thus the number of dislocations involved. With large grain sizes, the pile-ups will contain larger numbers of dislocations, which will in turn cause higher stress concentrations in neighboring grains.

In practical terms, the finer the grain size, the higher the resulting yield stress and, as a result, in modern steel working much attention is paid to the final ferrite grain size. While a coarse grain size of d^-1/2 = 2, i.e. d = 0.25 mm, gives a yield stress in mild steels of around 100 MPa, grain refinement to d^-1/2 = 20, i.e. d = 0.0025 mm, raises the yield stress to over 500 MPa, so that achieving grain sizes in the range 2-10 μm is extremely worthwhile.

Dispersion strengthening

In all steels there is normally more than one phase present, and indeed it is often the case that several phases can be recognized in the microstructure. The matrix, which is usually ferrite (bcc structure) or austenite (fcc structure) strengthened by grain size refinement and by solid solution additions, is further strengthened, often to a considerable degree, by controlling the dispersions of the other phases in the microstructure. The commonest other phases are carbides formed as a result of the low solubility of carbon in α-iron. In plain carbon steels this carbide is normally Fe₃C (cementite), which can occur, in a wide range of structures from coarse lamellar form (pearlite), to fine rod or spheroidal precipitates (tempered steels). In alloy steels, the same range of structures is encountered, except that in many cases iron carbide is replaced by other carbides, which are thermodynamically more stable. Other dispersed phases which are encountered include nitrides, intermetallic compounds, and, in cast irons, graphite.

Most dispersions lead to strengthening, but often they can have adverse effects on ductility and toughness. In fine dispersions (where ideally small spheres are randomly dispersed in a matrix) are well-defined relationships between the yield stress, or initial flow stress, and the parameters of the dispersion.

These relationships can be applied to simple dispersions sometimes found in steels, particularly after tempering, when, in plain carbon steels, the structure consists of spheroidal cementite particles in a ferritic matrix. However, they can provide approximations in less ideal cases, which are the rule in steels, where the dispersions vary over the range from fine rods and plates to irregular polyhedral. Perhaps the most familiar structure in steels is that of the eutectoid pearlite, usually a lamellar mixture of ferrite and cementite. This can be considered as an extreme form of dispersion of one phase in another, and undoubtedly provides a useful contribution to strengthening.

An overall view

Strength in steels arises from several phenomena, which usually contribute collectively to the observed mechanical properties. The heat treatment of steels is aimed at adjusting these contributions so that the required balance of mechanical properties is achieved. Fortunately the γ/α phase change allows great variations in microstructure to be produced, so that a wide range of mechanical properties can be obtained even in plain carbon steels. The additional use of metallic alloying elements, primarily as a result of their influence on the transformation, provides an even greater control over microstructure, with consequent benefits in the mechanical properties.

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Alloy steels are defined as those steels that:

1. contain manganese, silicon, or copper in quantities greater than the maximum limits (1.65% Mn, 0.60% Si, and 0.60% Cu) of carbon steel; or
2. that have specified ranges or minimums for one or more other alloying additions.

The low-alloy steels are those steels containing alloy elements, including carbon, up to a total alloy content of about 8.0%.

Except for plain carbon steels that are micro alloyed with just vanadium, niobium, and/or titanium, most low-alloy steels are suitable as engineering quenched and tempered steels and are generally heat treated for engineering use.

Low-alloy steels with suitable alloy compositions have greater hardenability than structural carbon steel and, thus, can provide high strength and good toughness in thicker sections by heat treatment. Their alloy contents may also provide improved heat and corrosion resistance. Effect of manganese on precipitation strengthening is greater than its effect in niobium steels. However, the absolute strength of niobium steel with 1.2% Mn is only about 50 MPa less than that of vanadium steel but at a much lower alloy level (that is, 0.06% Nb versus 0.14% V).

Another factor affecting the strength of vanadium steels is the ferrite grain size produced after cooling from-the austenitizing temperature. Finer ferrite grain sizes can be produced by either lower austenite-to-ferrite transformation temperatures or by the formation of finer austenite grain sizes prior to transformation.

The austenite grain size of hot-rolled steels is determined by the recrystallization and grain growth of austenite during rolling. Vanadium hot-rolled steels usually undergo conventional rolling but are also produced by recrystallization controlled roiling. With conventional rolling, vanadium steels provide moderate precipitation strengthening and relatively little strengthening from grain refinement. The maximum yield strength of conventionally hot-rolled vanadium steels with 0.25% C and 0.08% V is about 450 MPa. The practical limit of yield strengths for hot-rolled vanadium-microalloyed steel is about 415 MPa.

Niobium Microalloyed Steels. Like vanadium, niobium increases yield strength by precipitation hardening; the magnitude of the increase depends on the size and amount of precipitated niobium carbides. However, niobium is also a more effective grain refiner than vanadium. Thus, the combined effect of precipitation strengthening and ferrite grain refinement makes niobium a more effective strengthening agent than vanadium. The usual niobium addition is 0.02 to 0.04%, which is about one-third the optimum vanadium addition. Strengthening by niobium is 35 to 40 MPa per 0.01% addition.

Niobium steels are produced by controlled rolling, recrystalization controlled rolling, accelerating cooling, and direct quenching. The recrystallization controlled rolling of niobium steel can be effective without titanium, while recrystallization rolling of vanadium steels requires titanium for grain refinement.

Vanadium-Niobium Microalloyed Steels. Steels microalloyed with both niobium and vanadium provide higher yield strength in the conventionally hot-rolled condition than that achievable with either element alone. As conventionally hot rolled, the niobium-vanadium steels derive almost all of their increased strength from precipitation strengthening and therefore have high ductile-brittle transition temperatures. If the steel is controlled rolled, the addition of both niobium and vanadium together is especially advantageous for increasing the yield strength and lowering ductile-brittle transition temperatures by grain refinement.

Usually niobium-vanadium steels are made with relatively low carbon contents. This reduces the amount of pearlite and improves toughness, ductility, and weldability. These steels are frequently referred to as pearlite-reduced steels.

Niobium-Molybdenum Microalloyed Steels. Steels microalloyed with niobium and molybdenum may have either a ferrite-pearlite microstructure or an acicular ferrite microstructure. In ferrite-pearlite niobium steels, the addition of molybdenum increases the yield strength and tensile strength by about 20 MPa and 30 MPa, respectively, per 0.1% Mo, over a range of 0% to 0.27% Mo.

The principal effect of molybdenum on the microstructure is to alter the morphology of the pearlite and to introduce upper bainite as a partial replacement for pearlite. However, because the individual strength values of pearlite and bainite are somewhat similar, it has been proposed that the strength increase is due to solid-solution strengthening and enhanced precipitation strengthening caused by a molybdenum-niobium synergism.

Vanadium-Nitrogen Microalloyed Steels. Vanadium combines more strongly with nitrogen than niobium does, and forms VN precipitates in vanadium-nitrogen steel. Nitrogen additions to high-strength steels containing vanadium have become commercially important because the additions enhance precipitation hardening.

Some producers use nitrogen additions to assist in the precipitation strengthening of controlled-cooled sheet and plate with thicknesses above 9.5 mm. In one case, hot-rolled plates with vanadium and 0.018 to 0.022% N have been produced by controlled cooling in thicknesses up to 16 mm with yield strengths of 550 MPa. However, delayed cracking is a major problem in these steels. The use of nitrogen is not recommended for steels that will be welded because of its detrimental effect on notch toughness in the heat-affected zone.

Titanium-Microalloyed Steels. Titanium in low-carbon steels forms into a number of compounds that provide grain refinement, precipitation strengthening, and sulfide shape control. However, because titanium is also a strong deoxidizer, titanium can be used only in fully killed (aluminum deoxidized) steels so that titanium is available for forming into compounds other than titanium oxide. Commercially, steels precipitation strengthened with titanium are produced in thicknesses up to 9.5 mm in the minimum yield strength range from 345 to 550 MPa with controlled rolling required to maximize strengthening and improve toughness.

Like niobium and/or vanadium steels, titanium microalloyed steels are strengthened by mechanisms that involve a combination of grain refinement and precipitation strengthening; the combination depends on the amount of alloy additions and processing methods. In reheated or continuously cast steels, small amounts of titanium (<0.025% Ti) are effective grain refiners because austenite grain growth is retarded by titanium nitride.

Titanium-Niobium Microalloyed Steels. Although precipitation-strengthened titanium steels have limitations in terms of toughness and variability of mechanical properties, research has shown that an addition of titanium to low-carbon niobium steels results in an overall improvement in properties. Titanium increases the efficiency of niobium because it combines with the nitrogen-forming titanium nitrides, thus preventing the formation of niobium nitrides.

Acicular Ferrite (Low-Carbon Bainite) Steels. Another approach to the development of HSLA steels is to obtain a very fine, high-strength acicular ferrite microstructure, instead of the usual polygonal ferrite microstructure during the cooling transformation of ultra-low carbon (<0.08% C) steels with sufficient hardenability (by additions of manganese, molybdenum, and/ or boron). Niobium can also be used for precipitation strengthening and grain refinement. The principal difference between the structure of acicular ferrite (which is also referred to as low-carbon bainite) and that of polygonal ferrite is that the former is characterized by a high dislocation density and fine, highly elongated grains that are not exhibited in polygonal ferrite.

Acicular ferrite steels can be obtained by quenching or, preferably, by air-cooling with suitable alloys for hardenability. The principal advantage of this type of HSLA steel is the unusual combination of high yield strengths (415 to 690 MPa), high toughness, and good weldability. A major application of these steels is line pipe in arctic conditions.

For considering heat treatment of this group, several typical tool steels are selected as examples, designated only by the type letter and numeral as used in the USA and the UK for standardized tool steels, e.g. H13, O1. These designations are so well known by steel consumers all over the world that no qualifying institutional designations are necessary. Steels for which there are no AISI or BS specifications are designated according to Swedish (SIS) standards.

The hardenability of SIS 2550 is considerably good. SIS 2550 is air hardened in fairly heavy sections, which is of advantage where dimensional stability is concerned. Due to the lower carbon content the toughness is greater than the most of the other cold-work steels. When used for cold-work tools, the steel is tempered at 170-250°C, the resulting hardness then being 55-58 HRC. With regard to impact strength this steel, too, is susceptible to tempering treatments around 300°C (see Figure 1).

Figure 1. Steel SIS 2550. Hardness and impact strength as functions of tempering temperature

SIS 2550, after hardening and tempering at 200-250°C possesses very high tensile strength and good impact strength. The values given below have been obtained on tensile test specimens that were oil quenched from 830°C and tempered at 250°C.

Rp0,02 = 1370 MPa	A5 = 8%
Rp0,1 = 1520 MPa	Z = 33%
Rp0,2 = 1630 MPa	HRC = 54
Rm = 2000 MPa

Such favourable mechanical properties make the steel suitable for tools subjected to large static and dynamic forces. Some typical applications are dies for tableware, shear blades for heavy plate and dies for plastic moulds, which requires steel possessing a high degree of dimensional stability and excellent polishability.

SIS 2550 is also used for hot-work tools working at moderate temperatures, e.g. drop-forging dies. Such tools are tempered between 400°C and 600°C, the exact temperature depending on the hardness required and the working temperature of the tool. For working temperatures above approximately 400°C the hardness of the steel falls relatively quickly.

If higher working temperatures are involved it is recommended to use the special hot-work steels.

Grade S1 has both high wear resistance and high impact strength. The hardenability is inferior to that of the Cr-Ni-Mo steel SIS 2550. This implies that for dimensions greater than 50 mm in diameter, this steel is contour-hardening which, in fact, further increases its toughness.

The normal hardening temperature is about 900°C but it may be raised to 950°C without any risk of grain growth being incurred. If a hardness higher than 50 HRC is required in dimensions up to about 60 mm in diameter the steel should be quenched in oil. For heavier dimensions a combined water-oil quenching procedure may be necessary.

Of the many cold-work applications for tool steel, special mention should be made of the cold punching of plate having a thickness greater than about 3 mm. If a plate of increasing thickness is being punched and consequently the thickness measurement of the plate is approaching the diameter of the hole, the punches used show an increasing tendency to break if they are made from, for example, grades O1, A2 or D2. For this type of punching work, grade S1 has been shown to possess the best combination of toughness and wear resistance. A suitable hardness is 56-58 HRC.

Wear resistance is further increased if, during the course of the hardening treatment, the tools are heated for some 20 min in a cyanide bath. After this treatment no further finishing is required; at the most a very light finish grinding is permissible. Another example is the use of this steel as the impact hammer in nail guns, used for driving nails into concrete.

Owing to its high toughness in comparatively large dimensions, grade S1 can successfully be used for tableware dies, which, depending on their dimensions should either be quenched in oil or be heat treated according to the combined oil-water quenching procedure.

Another field of application is shear blades for cold shearing of heavy plate. Because of its rather good resistance to tempering, grade S1 may also be used for hot shears, a suitable hardness for this latter use being about 45 HRC.

In the field of hot-work, grade S1 has been superseded by other grades, e.g. H13. However, mention should be made of an interesting and successful field of application for grade S1 — as chisels used in process of electrolytic reduction of Aluminium from bauxite. The function of the chisels is to break up the hard alumina-containing crust which forms on the metal bath. During their use the chisels also come into contact with the bath itself and are thus subjected to both high temperatures and impact stresses. A suitable chisel hardness is about 350 HB.

Two steels have been chosen from this group as examples for the discussion, grade O1 (RT 1733) and Swedish SIS 2092 (SR 1855).

When carbon steel is used for punching dies or cold hobbing tools the dimensions of the tool are bound by a ruling section that is determined by the load on the tool. A punch or a die, made from carbon steel, having a diameter of, say, 50 mm, will show rather poor resistance to sinking on account of the shallow depth of hardening.

Should this resistance not suffice, another steel will have to be chosen, in this case grade O1 or SIS 2092. From the point of view of heat treatment, these two steels differ somewhat since their hardening temperatures are different. Steel SIS 2092 requires 850-890°C whereas grade O1 requires 800-840°C. Owing to its lower hardening temperature, O1 has somewhat greater dimensional stability. This property makes it a first choice for blanking dies and other tools requiring a high degree of dimensional stability (Figure 1).

Figure 1. Blanking tool made from steel O1

In both steels the depth of hardening decreases by roughly the same amount as the thickness of the section increases. In the diagram in Figure 2 the hardening temperature was raised as the cross-sectional area increased in order to increase the hardenability of the steel. Tools having diameters greater than about 80 mm or equivalent sections in flat dimensions are difficult to harden to full hardness if there are re-entrant corners. For such designs it is advisable to choose SIS 2092 since it obtains full surface hardness more readily and gives a more regular depth of hardening in a tool with varying section thickness.

Figure 2.

Curves showing depth of hardening for steel O1. Specimen 25 mm diameter oil quenched from 800°C. Specimen 50 mm diameter oil quenched from 820°C. Specimen 100 mm diameter oil quenched from 840°C

This point is well illustrated in Figure 3 which shows longitudinal sections through test specimens that have been hardened as normally prescribed for each grade concerned, i.e. oil quenching for both, from 820°C for grade O1 and from 870°C for SIS 2092.

Figure 3.

Longitudinal section (etched) through stepped test specimens made from: a) SIS 2092 and b) AISI O1. The diameters are 50, 75 and 100 mm

The above-cited example is to be regarded as a practical assertion of the possibility of estimating the depth of hardening from the Jominy diagram. However, it should be emphasized once again that a `contour-hardened` tool is tougher than a through-hardened one. Figure 4 shows a section through a `contour-hardened` Pilger roll.

Figure 4.

Transverse section through a Pilger roll made from SIS 2092. Size 50x120 mm

As a rule both steels are oil quenched. For heavy sections, e.g. dimensions greater than 100 mm in diameter, it is best to use water quenching when dealing with SIS 2092. When the surface temperature of the steel has fallen to between 400°C and 300°C the water quenching is interrupted by transferring the tool to an oil bath.

The tempering temperature for both steels is generally in the range 170-200°C which gives a hardness of more than 60 HRC. As can be seen from Figure 5, SIS 2092 has a greater resistance to tempering than grade O1.

On being tempered in the range 250-350°C the steel suffers a reduction in its impact strength, which in turn increases the risk of chipping. For this reason tools that are subjected to impact stresses should not be tempered in this temperature range. The higher impact strength manifested after tempering at 170-200°C is due to the presence of retained austenite, viz. about 10%.

Figure 5. Tempering curves for steel O1 and SIS 2092

This soft retained austenite can accommodate impact stresses better than the harder constituents. Retained austenite is decomposed when it is tempered at about 300°C.

If, during service, some areas of the tool have to support excessive pressures, for example the shearing edge of circular slitting knives (see Figure 6), retained austenite may be transformed to martensite, with spalling at the edge as a result. Should this occur, tempering at 300-400°C is recommended. After such a treatment the hardness of SIS 2092 still remains around 60 HRC.

Since the wear resistance of SIS 2092 is as much as 25% greater than that of grade O1, the former is very popular when a wear-resisting steel is required that can give a better performance than grade O1. Compared with this steel, SIS 2092 has been shown to have a considerably longer service life, particularly as drawing die steel.

Another interesting application is as cane-slitting tools (see Figure 7). The requirements of this type of tool are both high wear resistance and toughness in its thin walls. Of all the steels tested the best results were obtained with SIS 2092. In recent years SIS 2092 has increasingly been used for so-called Pilger rolls, which are in part made as rings and in part as dies.

Figure 6. Circular shears (slitting knives) made from SIS 2092

Figure 7. Cane-slitting tool made from SIS 2092

Figure 8 shows one of the world`s largest Pilger rolls, designed for cold-rolling 10 inches tubes. The only steel suitable for this tool was SIS 2092. Another field of application is for what are known as Yoder rolls. A sketch showing the principle of operation and tube manufacture is shown in Figure 9. For this mill unit, the wear resistance of rolls made from SIS 2092 has shown itself to be on a par with that of grade D2, in fact in some instances it has outlasted this grade.

Figure 8.

Pilger roll made from SIS 2092 for 10 in tube mill. Dimensions: 800 mm diameter x 400 mm, weight 800 kg

Figure 9.

Sketch showing tube-mill operating principle for welded tubes (Yoder mill). Welding stage omitted

This observation is particularly striking when stainless steel tubes are being rolled, since there is no `galling` when SIS 2092 is being used.

Reckoned on a tonnage basis, tool steel represents only a few percent of the total quantity of steel produced but its importance to the industry as a whole is immense. Regrettably this fact is seldom sufficiently appreciated. Perhaps in greatest measure this applies to the heat treatment of tool steel.

The cost of the steel and its heat treatment amounts generally to less than a quarter of the total cost of the whole tool. A wrong choice of steel or faulty heat treatment may give rise to serious disruption of production and higher costs.

In this text, an example is tool steel W1, designated only by the type letter and numeral as used in the USA and the UK for standardized tool steels. This designation system is so well known by steel consumers all over the world that no qualifying institutional designations are necessary.

Carbon steels and vanadium-alloyed steels

The hardening of these steels, which are made with carbon contents between 0,80% and 1,20%, is quite straightforward: Since the rate of carbide dissolution proceeds rapidly, the holding time, as a consequence, is short and therefore the heating of small tools can often take place without any extra precautions against atmospheric oxidation.

The hardening temperature is about 780°C. Quenching is carried out direct into brine with tempering following immediately. The quenching operation is the most critical part of the heat treatment since too slow a rate of cooling might give rise to either soft spots or quenching cracks.

If the tool is designed to contain hardened areas around holes or reentrant angles the cooling effect must be very intensive at these areas. Manual stirring will often suffice but in many cases the coolant must be sprayed on to the tool. For sections heavier than 20 mm the depth of hardening, i.e. the distance from the surface to the 550 HV level, is about 4 mm. Sections less than about 8 mm in thickness will harden through.

For awkward tools, hardenability may be a crucial factor and under such circumstances the composition of the steel must be adjusted in accordance herewith, in particular as regards the alloying elements Mn and Cr, which have a powerful influence on hardenability.

The diagram in Figure 1 shows how the hardening temperature affects the depth of hardening and fracture number on Wl-type steel of conventional composition. The V-content is only 0,04%, which implies that the steel starts to be coarse-grained when the hardening temperature exceeds 815°C.

Figure 1.

Depth of hardening for carbon steel, 25 mm in diameter, corresponding to W1. Quenched in water from various temperatures

In Figure 2 are shown the results of corresponding trials with steel containing somewhat larger amounts of alloying elements. The depth of hardening is considerably greater. Owing to the high content of V the steel remains fine-grained even when hardened from exceptionally high temperatures.

The very considerable toughness inherent in plain-carbon steel, due to its shallow-hardening properties, is forfeited if the tool through-hardens locally at some sections because the cross-sectional area there is too small. For shearing tools or small tools generally, such as scissors, knives or letter die punches, which are not subjected to heavy impact blows, this drawback is of less importance. Tools operating under heavy blows, e.g. upsetting dies for cold-heading of bolts, must not be through-hardened.

Coining and striking punches are other examples of carbon tool steels that require high wear resistance. Such tools may also be subjected to bending stresses and should therefore not be through-hardened. The tempering temperature normally used for tools belonging to this group lies in the range 170°C, the hardness being generally about 60-64 HRC. Representative examples of tools made from grade W1 are shown in Figure 3.

Figure 2.

Depth of hardening for carbon steel, 25 mm in diameter, corresponding to W1. Quenched in water from various temperatures

Figure 3. Punches made from steel W1

Wire ropes for haulage purposes are usually made from carbon steel wires ranging from 0,35 to 0,5% carbon, and before drawing the material is subject to a heat-treatment known as patenting.

Patenting consists of passing the wire through tubes in a furnace at about 970^oC. This high temperature treatment produces uniform austenite of rather large grain size. The subsequent cooling - in air or molten lead - is rapid since the sections treated are generally small (e.g. wire rods), so that the resulting structure consists of very fine pearlite preferably with no separation of primary ferrite.

The large crystals would give rise to brittleness if the material was left in the heat-treated condition, but this effect is not noticed after a few drawing passes. Variation in hardness - either softer or harder - can be produced by tempering martensite, but such material does not draw so well as patented wire, which is able to withstand reductions of area up to 90%. The strength is explained on the basis of the reduced ferrite cells and the alignment of cementite in fibres.

Hardenability, mass effect, ruling section

Hardenability is the measure of the depth to which a steel will harden on quenching.

The maximum hardness is mainly a function of the carbon content. The hardenability of steel depends on:

1. The quenching medium and method of quenching.
2. Composition of the steel and method of manufacture.
3. Section of the steel.

The so-called "Mass effect" arises from the fact that even with the most severe quench the cooling of a bar is progressively slower, from the outside to the centre due to the low thermal conductivity of the steel. It must be appreciated, therefore, that it is the rate of cooling of a piece of steel which determines the properties resulting from a quenching process, and not mass or weight.

The effects of mass are shown by the following results (Table 1) on quenched bars of varying section, from the centre of which standard test specimens were machined. To prevent the misuse of steels it has become desirable to state the limit of thickness or "ruling section" up to which the mechanical properties quoted can be obtained. For sections other than cylinders, reference should be made to BS 970 for conversion factors.

One of the most important functions of alloying elements in high tensile steels is to produce fully hardened structures in large sections.

Table 1.

Effect of mass on the properties of carbon steel: C-0,45%, Si-0,32%; Mn-0,78%; S and P-0,02%

Dia. of bar, mm	Tensile Strength Rm, MPa	Elongation, A, %	Izod, J
Water quenched from 870oC and not tempered
17	1853	3	4
29	1035	8	18
76	803	15	32
Tempered at 600oC
17	850	18	103
76	741	25	59

Shallow hardening steels (C 1%; Mn 0,3%) are often used for dies in which a tough core and a hard skin is required. The depth of hardening is increased by drastic quenching media, and by manganese and elements such as nickel and chromium, but decreased by fine austenite grains.

The addition of 3% nickel and 1% chromium retards the transformation of austenite so much that oil quenching exceeds the critical cooling velocity and all portions from edge to centre of average sections will harden. The results in Table 2 illustrate this effect and should be compared with the ones in Table 1.

Table 2.

Effect of mass on properties of alloy steel: C-0,31%, Si-0,14%; Mn-0,70%; Ni-3,27%, Cr-0,82%

Dia of bar, mm	Tensile Strength Rm, MPa	Elongation, A, %	Izod, J
Oil quenched from 820°C, tempered at 600°C
17	958	24	85
76	927	22	82

It has now been found that small additions of a number of alloying elements are more effective than an equivalent total addition of one element, and enables appropriate use to be made of alloy scrap.

In the Jominy test (BS 4437) a 25 mm dia bar 100 mm long is heated to the normal hardening temperature, then inserted in a standard jig and a 12,5 mm dia jet of water at 24^oC is directed against one end of the test piece. The free height of the jet is 62,5 mm.

When cold, hardness measurements are made along flats on the bar and these are plotted against distance from the quenched end as in Fig. 1, which also includes the cooling rates at various distances from the quenched end. The positions along the Jominy bar having cooling rates equal to those at the centres of bars of various diameters are marked in Fig. 1 for water and oil quenching.

Figure 1. Hardness distance curves for end quench test samples

From the Fig. 1 the effectiveness of a given quench on various steels can be judged; for example, bars of 1,5% Ni Cr Mo steel up to 100 mm dia can be satisfactorily hardened in oil, whereas for the plain 0,95% carbon steel the maximum diameter of the fully hardened bar will be less than 12,5 mm for the same quench.

The Jominy end quench test is not only a very reproducible test. but is also an excellent laboratory test for classifying steels in order of merit as regards hardenability and it gives a very useful first approximation to the merits of a given steel. As to whether a steel with a particular hardenability will give the satisfactory mechanical properties in a particular section must be decided by trial.

Actually, the method of trial - apart from giving the final answer - is also, in many cases, far simpler than the involved mathematical processes which have been developed in an endeavour to predict the mechanical properties of a steel in any given section from the Jominy curve.

Figure 2. Designation of hardenability limits

The hardenability of different casts of steel of the same analysis may vary substantially and similarly significant variations in hardenability may occur in samples taken from different parts of the same ingot. As a consequence Jominy results are given as bands not simple curves and BS 4437 uses two points to specify hardenability as shown in Fig. 2 and as follows:

1)	A desired hardness value at 2 desired distances, e.g.	A-A
2)	Minimum and maximum hardness values at desired distance	B-B
3)	Two maximum hardness values at 2 desired distance	C-C
4)	Minimum hardness values at 2 desired distances	D-D
5)	Any minimum and any maximum hardness.