Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

12.16 Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels AB Nissan and KO Findley, Colorado School of Mines, Golden, CO, USA Ó 2014 Elsevier Ltd. All rights reserved.

12.16.1 Fundamentals of Induction Hardening 12.16.1.1 Overview of the Induction-Hardening Process 12.16.1.2 Hysteresis Losses, Eddy Currents, and Heat Transfer 12.16.1.3 Alloy Content and Carbon Content Selection of Induction-Hardened Steel 12.16.1.4 Review of Critical Induction-Hardening Parameters 12.16.2 Rapid Austenitization and Hardenability of Steel Microstructures 12.16.3 Residual Stress and Its Effects on Mechanical Properties 12.16.4 Characterization of Induction-Hardened Components 12.16.5 Strength and Fatigue of Induction-Hardened Parts 12.16.5.1 Estimation of Fatigue Crack Nucleation Location 12.16.5.2 Fatigue Fracture Surface Location and Morphology 12.16.6 Industrial Standards Relevant to Induction Hardening 12.16.7 Summary and Conclusions Acknowledgments References

12.16.1

581 581 583 585 586 587 590 592 595 598 600 602 603 603 603

Fundamentals of Induction Hardening

Induction hardening is quick and cost effective, and allows for repeatable high throughput of samples with a very reliable low maintenance setup, and it offers significant advantages over other surface-hardening methods. The short time between heating and quenching does not allow sufficient time for decarburization or significant grain growth of the steel, and there is little to no distortion of the part from the heat-treatment operation. Axisymmetric or near axisymmetric parts are ideal for induction hardening, but more complex geometries such as gears can also be induction hardened. Induction hardening is different from other case-hardening methods in that there is no chemical change at the surface of the part after induction hardening, unlike surface-hardening treatments such as carburizing or nitriding. Instead, induction hardening relies exclusively on phase transformations to create a wear-resistant case and compressive residual stresses at the surface. The resulting martensitic case as well as the compressive residual stresses improves the wear and fatigue life of induction-hardened components significantly over nonhardened components. Any metal can be induction hardened, but ferromagnetic materials (e.g., iron or steel) are especially responsive to induction hardening due to their ferromagnetic properties. Ferromagnetic materials respond well to induction heating because both hysteresis and eddy current losses can contribute to heating in a part, depending on the frequency of the induction-hardening operation (Figure 1). Low-alloy steel can be effectively induction hardened to produce substantial improvements in wear and fatigue performance without the addition of expensive alloying elements, and induction-hardening equipment is also very versatile in that it can be used for either case or through hardening heat treatments as well as selectively hardening critical regions.

12.16.1.1 Overview of the Induction-Hardening Process The basic induction-hardening setup employs a water-cooled copper conductor that surrounds the workpiece. Copper coils are used to minimize resistive heating from the high alternating current, and the coil is water cooled to prevent melting due to the highpower input. Several examples of induction coil/workpiece configurations are shown in Figure 2. Induction-hardening configuration design is also versatile; for example, a single surface or just the inner diameter of a tube can be heated. Induction hardening can be conducted with scanning coils or multiturn coils. Either single-shot (stationary) or scanning (progressive) induction-hardening processes are possible with the water-cooled copper coil configuration. Single-shot induction hardening is usually conducted with a multiple-turn coil where the coil and sample do not move relative to each other (Figure 2(a)). Scanning induction heating involves holding the sample stationary and moving the coil up or down the workpiece at a constant scan speed (Figure 2(b)). The main advantage of a scanning coil is that the power input to the workpiece is concentrated into a single coil, which can allow for faster heating rates and shallower case depths depending on the size of the power source available for hardening. Additionally, the scanning induction coils can have either an integrated quench ring (Figure 3(a)), which is incorporated into the coil and allows for immediate quenching of the material after heating, or an external quench ring separate from the scanning coil (Figure 3(b)), which allows for a greater flow rate of quenchant on the sample. Single-shot coils do not have integrated quench rings and instead rely on an external quench apparatus. The workpiece is constantly rotated during both single-shot and scanning induction of axisymmetric geometries to ensure even heating.

Comprehensive Materials Processing, Volume 12

http://dx.doi.org/10.1016/B978-0-08-096532-1.01218-8

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Heat producing losses

Eddy current losses

Hysteresis losses

Frequency (kHz) Figure 1 Heat contribution from eddy currents and hysteresis losses below the Curie temperature as a function of frequency. Adapted from Curtis, F. W. High Frequency Induction Heating; Lindsay Publications Inc.: Bradley, IL, USA, 1987.

Induced eddy currents

Magnetic field

Inclusion coil

OD heating single-turn coil

OD heating multiturn coil Magnetic field

Workpiece

(a)

ID heating multiturn coil

Magnetic field

Inductor

(b)

Magnetic field

Surface heating single-turn coil

Figure 2 Typical induction coil to workpiece configuration for induction heating of (a) outer and inner diameter heating with multiturn coils and (b) outer diameter and surface heating with single turn coils. Reproduced from the Haimbaugh, R. E. Practical Induction Heat Treating; ASM International: Materials Park, OH, USA, 2001.

Figure 3

Scanning coil with (a) an integrated quench ring and (b) a separate quench ring.

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Figure 4 Typical frequency and power settings for different types of induction heating applications. Adapted from Semiatin, S. L.; Stutz, D. E. Induction Heat Treatment of Steel; American Society for Metals: Metals Park, OH, 1986.

The induction-hardening generator power and frequency specification varies by manufacturer and model but can range from 10 kW to 1 MW using alternating current at frequencies of 0.1 kHz to greater than 200 kHz. The power and frequency used depend on the desired case depth and on whether scanning or single-shot induction hardening is utilized (2–5) (Figure 4). As shown in Figure 4, single-shot hardening is usually conducted using higher power than scan hardening because the entire section to be hardened must be heated at once rather than progressively, which requires greater power. The frequency range at which scan hardening is conducted is also more variable than single-shot hardening, and the power input requirement decreases as the scan hardening frequency increases. During scan hardening, the penetration depth of magnetic induction decreases as the frequency is increased, which results in increased power densities (assuming power remains constant). Because of the increase in power density, less time is required for heating.

12.16.1.2 Hysteresis Losses, Eddy Currents, and Heat Transfer Induction hardening is a complex interaction of both electromagnetic radiation and heat transfer. During the heating cycle, a ferromagnetic material (e.g., steel) is heated through both hysteresis and eddy current losses. Hysteresis losses dominate at frequencies less than approximately 60–70 kHz (1) and are caused by the rapid change of magnetic fields due to the alternating current in the induction coil. The changing magnetic field causes the domains in the ferromagnetic steel to rapidly change, which results in a hysteresis in the magnetization versus magnetic field strength behavior as shown in Figure 5. The magnitude of the power dissipated into the workpiece is the area encompassed within the hysteresis curve (1–7). As the frequency of the alternating

Magnetic induction (B)

Applied magnetic field (H)

Figure 5

Typical hysteresis curve for a ferromagnetic material where B is the magnetic induction and H is the externally applied magnetic field.

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

current of the induction coil is increased, the hysteresis curve contracts until the forward and reverse responses are the same, and thus the power input from hysteresis losses decreases and is surpassed by eddy current losses. Eddy current heating, also known as joule heating, is the primary form of heat input into a workpiece at frequencies greater than 70 kHz (1). Eddy currents are caused by local magnetization of the workpiece, which causes magnetic eddies to develop within the material. The eddies dissipate their energy into the surrounding material, resulting in heating of the workpiece. For thin sheets, assuming the magnetic field is uniform and the material is homogeneous, the power input into a workpiece due to eddy current losses is (5): P¼

p2 B2v d2 f 2 6rD

[1]

where Bv is the peak magnetic flux density, d is the thickness of the sheet, f is the frequency, r is the resistivity of the material, and D is the density of material. Thus, the power input into a workpiece increases with increasing frequency and magnetic field strength and decreases as the resistivity of the material increases. Since the magnetic field during induction hardening of a ferromagnetic material is not uniform throughout the thickness of the part, the depth at which the majority of current and power are concentrated is referred to as the penetration depth or skin effect. The penetration depth, d (m), is the distance from the surface into the workpiece at which the current is 37% and the power density is 14% of the current and power at the surface of the part (63% of the current and 86% of the power are contained in the penetration depth) (7). It can be predicted by (8): 1 d ¼ pffiffiffiffiffiffiffiffiffiffiffiffiffi pf mr s

[2]

where f is the frequency (Hz), mr is the relative magnetic permeability (unitless), and s is the electrical conductivity (S m1). The relative magnetic permeability is a ratio of the magnetic permeability (m) of the material over the permittivity of free space (mo) and is given by (9): m [3] mr ¼ mo Equation [3] can be approximated by: 500 d ¼ pffiffiffi f

[4]

where d is the penetration depth (mm) and f is the frequency (Hz); the equation may be applied for temperatures greater than 800 C (10). As a general rule of thumb for induction hardening of steel components, the penetration depth should be twice the desired case depth (3). The induction-heating response of a ferromagnetic workpiece changes once the material passes through the Curie temperature, the temperature at which the material switches from ferromagnetic to paramagnetic behavior. In steel, the Curie temperature is 770  C regardless of carbon content. Above the Curie temperature, the relative magnetic permeability becomes unity, and the

Percentage of maximum penetration depth, δ

100

100

90

90 Maximum penetration depth, δ

80

80

70

70

60

60

50

50

40

40

30

30 Minimum penetration depth at start of heating (4% of maximum)

20

Curie temperature, 770 °C

10

20 10

0 0

200

400

600 800 Temperature, °C

1000

0 1200

Figure 6 Penetration depth, d, as a function of temperature for a plain carbon steel demonstrating the dramatic increase in penetration depth above the Curie temperature.

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

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Heat loss due to conduction

Heat loss due to radiation and convection

Induction coil

Induction coil

Quenchant

Case

Core

Figure 7 Scanning coil induction-hardening schematic demonstrating heat input from the induction coil as well as the heat losses from radiation, convection, and conduction.

penetration depth increases dramatically as shown in Figure 6, which plots the relative penetration depth compared to the maximum possible as a function of temperature. This increase in the penetration depth of the workpiece can result in overtempering of the core microstructure during hardening and longer heating times required for deep case depths because of more diffuse heating through the austenitized layer. Adding to the complexity of induction hardening is that all three types of heat transfer occur during processing: conduction, convection, and radiation (Figure 7). Heat is constantly conducted from the hot surface to the cooler core, so additional heat input is required to achieve the desired temperature (such as the austenitization temperature). Furthermore, heat conduction into the interior can overtemper the microstructure at the case/core interface. Heat is also lost to the environment through radiation and, to a lesser extent, through convection. Heat loss is a major factor in determining the total heat input into the workpiece. During quenching, convection draws heat from the surface, and conduction draws heat to the surface. However, heat also conducts to the core of the workpiece until the surface of the workpiece is cooler than the core, at which point heat is then conducted from the warmer core to the cooler surface. The heat conduction processes can result in a small amount of auto-tempering of the workpiece in cases where the workpiece and case depth are sufficiently large.

12.16.1.3 Alloy Content and Carbon Content Selection of Induction-Hardened Steel As previously stated, induction hardening, unlike other case-hardening methods, relies exclusively on the formation of martensite, with no composition change to produce a hard fatigue-resistant case and compressive residual stresses at the surface of the part. Increasing the strength of the case also increases the load-carrying capacity of the part in torsion. Because induction hardening relies on transformation of the case to martensite, carbon and alloy steels in the carbon content range of 0.4–0.5 wt. pct. C are especially suited for induction hardening to match the surface hardness expected from other surface-hardening processes (Figure 8). This carbon range allows for the high hardness (55–60 HRC) required for wear and fatigue resistance, but generally avoids the problems associated with higher carbon content steel such as quench cracking and quench embrittlement. Figure 9 shows that steels with carbon content greater than 0.5 wt. pct are susceptible to quench embrittlement, where the part becomes susceptible to brittle intergranular fracture due to cementite formation and phosphorus at prior austenite grain boundaries. However, carbon and alloy steels with carbon content outside of the 0.4–0.5 wt. pct. range are commonly used in industrial induction-hardening processes. If lower carbon steel is used, then a lower surface hardness and resulting wear resistance must be accepted. If a higher carbon content is used (C > 0.55 wt. pct.), precautions should be taken to ensure that quench cracking and quench embrittlement at the surface do not occur. Plain carbon steels (i.e., 10xx series) have a limited amount of hardenability. Hardenability is the depth or maximum diameter where a 50% martensitic microstructure can be achieved; it is expressed as an ideal diameter and can be calculated for any alloy content using ASTM A255-10 ‘Standard Test Methods for Determining the Hardenability of Steel’ (8). The use of alloying elements such as Ni, Mo, Mn, and Cr increases the hardenability of steel. By increasing the hardenability of an alloy, greater case depths can be achieved. In addition to alloying to improve hardenability, microalloy additions of Nb, V, Al, and Ti can be used to form carbides and nitrides that retard grain growth during the austenitizing process and improve the tempering response of the material, which may improve mechanical performance (11,12).

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

70

Typical surface hardness range for carburized components

Rockwell hardness C scale (HRC)

60

100% Martensite hardness

50

40 50% Martensite hardness 30

20 0.1

0.2

0.3

0.4 0.5 Carbon content (wt%)

0.6

0.7

Figure 8 Hardness range for steel with different carbon contents as well as the typical hardness range for carburized components (reproduced from the ASTM Standard A255-10, Standard Test Methods for Determining Hardenability of Steel; ASTM International: West Conshohocken, PA, USA, reprodued from the Abbaschian, R.; Abbaschian, L.; Reed-Hill, R. R. Physical Metallurgy Principles, 4th ed.; Cengage Learning: Stamford, CT, 2009, Reprodued From the Kalpakjian, S.; Schmid, S. R. Manufacturing Engineering and Technology; Prentice Hall: Upper Saddle River, NJ, 2001).

700 350 600 Tempered-martensite embrittlement region Carbide-induced reduced toughness

500

250

LTT martensitic steels

200 Ductile fracture region 150

Carbon-dependent toughness maxima

Quenchembrittlement region Intergranular fracture

100

High-carbon LTT steels used: (a) When intercritically austenitized (b) With surface compressive stress (c) Under compressive or hertzian loading

400

300

Tempering temperature (°F)

Tempering temperature (°C)

300

200

As-quenched low toughness region

50

100 0

0.2

0.4

0.6

0.8

1

Carbon content (pct) Figure 9 Schematic showing the ductile fracture region and quench-embrittlement regions as a function of tempering temperature and carbon content (reproduced from the Krauss, G. Steels: Processing, Structure and Performance; ASM International: Materials Park, OH, USA, 2005.).

12.16.1.4 Review of Critical Induction-Hardening Parameters The main parameters, excluding workpiece geometry, that determine the depth and speed of heating of a given workpiece are the alternating current frequency, alternating current power, scan speed (scanning coil), dwell time (single shot coil), type of quenchant,

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

587

and coil configuration. During processing, frequency, power, and scan speed/dwell time are the quickest and least costly to change to achieve the desired case depth. However, depending on the induction-heating system frequency, power or scan speed may be fixed. Frequency: As shown in eqns [2] and [4], frequency adjustment allows for changes in the penetration depth, d, of the workpiece. For relatively thin workpieces, a higher frequency must be used to prevent through hardening of the material. The frequencies of induction-hardening machines are generally broken into three categories: low frequency (<10 kHz), medium frequency (10–70 kHz), and high frequency (>70 kHz) (7). The high-frequency range is also termed ‘radio frequency.’ Medium- and highfrequency are used the most often for surface hardening of steel components. Power: A range of power settings can be used for induction hardening, from as low as 10 kW to over 1 MW of power. If, for example, more power is imparted to the workpiece, the heating rates within the skin depth region at a given frequency will be increased. Scan speed/dwell time: If the frequency and power are held constant, changing the scan speed/dwell time can change the temperature distribution and resulting case depth within the workpiece. Increasing the scan speed may be required to prevent melting of the sample and to decrease the case depth. Lowering the scan speed may be desired if the microstructure in the workpiece does not transform readily to austenite upon heating and requires greater heat input (e.g., spheroidized microstructure). Quenchant: A wide variety of quenchants are available in liquid (water, oil, or polymer) and gas (N, He, Ar, etc) form. Each of these media provides different levels of heat extraction at the surface of the part. Care must be taken to select a quenchant that is severe enough to form martensite at the desired case depth but not too severe as to cause quench cracking at the surface of the workpiece. For alloys with a carbon content and/or hardenability, a less severe quenchant should be utilized. Polymer-based quenchants are popular in induction-hardening applications because the percentage of polymer can be tuned to adjust the severity of the quench to the necessary level. Coil configuration: The design of induction coils is outside the scope of this chapter, but excellent descriptions on coil design, configuration, and troubleshooting of induction coils can be found in Ref. (2,3,7). Tempering conditions: The final hardness of the induction-hardened part after processing is a function not only of carbon content but also of the use of tempering after processing. A typical tempering operation after induction hardening is 176 C for one to one and a half hours. If a higher tempering temperature is used, then the surface hardness of the part will decrease, but the toughness of the case region will also increase. Because the highest possible wear and fatigue resistance in the case is generally desired, the 176 C temper is often utilized. It should be noted that independent of the tempering temperature, all inductionhardened parts should be tempered within 4 h of induction hardening in accordance with the Society of Automotive Engineers standard SAE AMS2745 (16).

12.16.2

Rapid Austenitization and Hardenability of Steel Microstructures

In order to form a martensitic case microstructure during induction hardening, the case microstructure must first be transformed to austenite before quenching. Induction hardening is a unique heat treatment in that the heating cycle is very short, so the extent of austenitization upon heating and at the brief hold at peak temperature is strongly dependent on the starting microstructure. Austenite can nucleate at ferrite grain boundaries, adjacent to spheroidized carbides, or within pearlite colonies; Figure 10 schematically shows nucleation of the austenite grains during heating at these various sites. Since austenite has a much higher solubility for carbon, nucleation near carbon-rich regions such as cementite is favorable. Growth of the austenite grain into the pearlite is very rapid due to the abundance of carbon from the cementite and the correspondingly short diffusion distances between carbon-rich regions in the interlamellar ferrite. Figure 11 demonstrates how austenite can nucleate at pearlite colony boundaries or at cementite lamellae at an initial time t1; the austenite grains are denoted A1, A2, and A3. Due to the slow dissolution of carbides, multiple austenite grains can nucleate within the same pearlite colony. The schematic at t2 shows that austenite growth is more rapid at the pearlite colony boundary (17). If, however, the austenite is to grow into pro-euctectoid ferrite, carbon must diffuse from the cementite, which results in much slower austenite growth.

A2

A1 A1

1 2 A2

2

A2

A1 1 2

3

3 Austenite

A3

1

Carbide A3 (a) Ferrite

(b) Spheroidite

(c) Pearlite

Figure 10 Austenite nucleation sites in (a) ferrite, (b) spheroidite, and (c) pearlite (reproduced from the Krauss, G. Steels: Processing, Structure and Performance; ASM International: Materials Park, OH, USA, 2005).

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 11 Austenite nucleation at the start of the nucleation and growth process in a pearlite colony (left schematic), and later within the austenite grain at remaining carbides (reproduced from the Speich, G. R.; Szirmae, A.; Richards, M. J. Formation of Austenite from Ferrite and Ferrite-Carbide Aggregates. Trans. Metall. Soc. AIME 1969, 245, 1063–1074).

Table 1 Prior microstructure, maximum surface hardness, and case depth for induction-hardened 1541 and 4140 alloys in which scan speed, power, and frequency were held constant and only prior microstructure and alloy were changed (18)

Prior microstructure

Alloy

Max. hardness after processing, HRC

Case depth (50 HRC), mm

Surface axial residual stress, MPa

Spheroidized

1541 4140 1541 4140 1541 4140 1541 4140 1541 4140 1541 4140

55–56 52–53 56–57 57–58 56–57 56–57 56–57 57–58 58–59 58–59 58–59 58–59

1.42 0.41 2.03 1.42 2.13 1.02 1.73 1.32 2.74 2.84 3.35 2.80

550 550 350 420 400 490 240 290 350 290 260 310

Ferrite þ spheroidized pearlite Ferrite þ coarse pearlite (small PAGS) Ferrite þ coarse pearlite (large PAGS) Quenched and tempered martensite As-quenched martensite

Because diffusion is sluggish compared to the short heating time during induction hardening, an ideal microstructure would have evenly distributed carbon or a fine dispersion of carbides throughout the microstructure to allow for rapid austenitization. Table 1, taken from the work of Coryell (18), provides example results for the effects of alloy type and starting microstructure on the induction-hardening response. Coarse spheroidized microstructures are the slowest and least responsive to induction hardening, and as-quenched martensite responds the best to induction hardening followed by quenched and tempered martensite (18). Carbon is supersaturated and relatively well dispersed in as-quenched martensite and forms fine carbides upon tempering that are readily dissolved during heat treatment. Table 1 also shows that the alloy with the higher hardenability, 1541, has a higher case depth for every condition. If the carbides within an austenite grain do not entirely dissolve during induction processing, the austenite is considered inhomogeneous. The lower carbon content of the austenite reduces the propensity for quench cracking and quench embrittlement for steels with carbon contents greater than 0.5 wt. pct. (15). However, if too many carbides remain in solution, the surface hardness of the induction-hardened part may be lowered. Because fully homogeneous austenite would require an almost infinite amount of time to form, austenite is considered homogeneous even if very small carbides are present. Besides microstructure, the extent of austenitization depends on heating rate and induction process temperature. During furnace heat treating, an equilibrium Fe–C phase diagram can generally be used to determine the critical temperatures for the ferrite þ austenite, Ac1, and austenite, Ac3, phase fields because the time required to reach these temperatures in a furnace are relatively sluggish. However, heating rates can be in excess of 1000  C s1 during induction hardening, so the influence of transformation kinetics on the critical transformation temperatures must be considered. At these high heating rates and short hold times at peak temperature, diffusion is very limited due to the short time at temperature. The Ac3 temperature can increase by 100–300  C from the equilibrium Ac3 temperature. The effect of heating rate is illustrated in Figure 12, which plots the Ac3 temperature as a function of heating rate in 1042 steel for three different starting heat-treatment conditions: annealed, normalized, and quenched and tempered (6). The Ac3 temperature increases with heating rate for all three conditions, and as described previously, the microstructure with the most homogeneous distribution of carbon, the quenched and tempered condition, transforms more readily

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

589

Figure 12 Differences in Ac3 temperature for a 1042 steel alloy with three different starting microstructures at different heating rates compared to the Ac3 temperature from the Fe–C phase diagram. Replotted from Semiatin, S. L.; Stutz, D. E. Induction Heat Treatment of Steel; American Society for Metals: Metals Park, OH, 1986.

Figure 13 Effect of heating rate and alloy composition on the A3 temperature. Replotted from Orlich, J.; Rose, A.; Wiest, P. Atlas zue Wa¨rmebehandlung der Sta¨hle – Band 3, Zeit-Temperatur-Austenitisierung-Schaubilder; Verlag Stahliesen: Düsseldorf, Germany, M.B.G.,1973.

(at lower temperatures) to austenite. Similarly, a time–temperature-austenitizing diagram is plotted in Figure 13 for three mediumcarbon steels that were induction heated at several heating rates. The critical temperature increases with increasing heat rate for all three alloys. During induction hardening, the heating rates are high and the quench rates can be extremely high, so induction-hardened components undergo a phenomenon termed superhardness (15). This phenomenon is associated with refinement of the austenite grains during heating due to the very short time at austenitizing temperatures. The newly formed fine-grained austenite is then quenched very rapidly to form very hard martensite. Induction processing can result in a 1–5 HRC increase in hardness compared to the martensite hardness expected for a given carbon content (Figure 14).

590

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 14 Hardness as a function of carbon content at the surface of induction-hardened plain carbon steels compared to the same steels subjected to furnace heat treatments (HT) and water quenching (WQ). Superhardness is observed at the surface of induction-hardened components compared to the HT and WQ steels. Adapted from Hassell, P. A.; Ross, N. V. Induction Heat Treating of Steel. In ASM Handbook; Vol. 4, Heat Treating, 1991.

12.16.3

Residual Stress and Its Effects on Mechanical Properties

Residual stresses can be beneficial or detrimental, depending on the residual stress profiles in the material after final processing. If the residual stresses at the surface are tensile, they are considered detrimental to the fatigue and most other mechanical properties of a component. If, however, the residual stresses at the surface of the part are compressive, the residual stresses are considered beneficial. Induction hardening generally produces a favorable residual stress profile that is compressive at the surface and tensile in the core (21). Compressive stresses arise at the surface of induction-hardened parts for the following reasons: 1. Residual stresses occur after induction processing primarily due to the volume difference between austenite, ferrite, and martensite (Figure 15). The larger the volume change upon the decomposition of austenite (e.g., formation of martensite, pearlite, etc.), the greater the tensile or compressive residual stresses. Figure 16 shows a schematic cooling profile of an

Figure 15 Specific volume of steel phases as a function of carbon content for a tool steel (reproduced from the Totten, G., Ed. Steel Heat Treatment Handbook; Marcel Dekker, Inc.: New York, USA, 1997).

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Figure 16 Example (a) CCT diagram demonstrating the cooling profile from the surface and the core of the samples and (b) the resulting residual stress profile after quenching.

induction-hardened part superimposed on a continuous cooling transformation (CCT) diagram. The outside (case) of a part transforms to austenite during induction processing, but the majority of the core microstructure is not heated to a temperature that allows for partial or full austenitization. Upon cooling, the surface of the part undergoes a phase transformation to martensite, which also involves volume expansion, while the core does not undergo a phase transformation. Thus, the volume expansion due to the phase transformation at the surface of the part results in residual compressive stresses, and the core is subjected to residual tensile stresses. Phase transformations have the largest effect on the residual stress distribution during induction processing. 2. During quenching, the surface of the part cools more rapidly and thus contracts more than the core, which initially results in residual tensile stresses at the surface and compressive stresses in the core. However, upon further cooling, the core begins to cool more rapidly and contracts more than the surface, and the residual stress profile flips so that the surface goes into compression and the core is under tension. Many techniques such as X-ray diffraction (XRD) (22,23), neutron diffraction (24), the contour method (25), and the slitting (crack compliance) method (26) have been developed to determine residual stresses in a component. Of these four techniques for residual stress determination, XRD is the most common. The depth profile is accomplished by chemical polishing or machining with chemical polishing into the surface of the sample and retaking measurements at each target depth. Residual stress is calculated from XRD results based on the deviation of the crystal lattice from the unstressed state. The interatomic spacing (d) of the lattice increases due to tensile residual stresses and decreases due to compressive stresses. The d-spacing is calculated from Bragg’s law (9): l ¼ 2d sin Q

[5]

where l is the wavelength of the characteristic X-ray radiation, d is the interplanar spacing, and Q is the angle between the incident beam (X-ray source) and the horizontal axis or the angle between the diffracted beam and the horizontal axis. Although Bragg’s law can be used to obtain residual stress values, the preferred method is the sin2 j method (23). This method utilizes a single theta angle and relies on rotation of the sample relative to the X-ray source and detector. Using a Cr–Ka target, the ideal 2-theta angle for ferrous alloys is 156 (d-spacing of 0.117 nm (23)). Several j angles, the rotation angle of the specimen relative to the X-ray beam (Figure 17), are then selected such as 0, 20, 29, and 36.3, which corresponds with sin2 j values of 0, 0.1170, 0.2340, and 0.3505, respectively (23). These angles are considered ideal because they are uniformly spaced as sin2 j values.

Figure 17 Schematic showing the basic geometry for stress measurements by XRD. Adapted from Residual Stress Measurement by X-ray Diffraction; 2003 ed., SAE International: Warrendale, PA, USA, 2003.

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 18 Sin2 j method of determining residual stress. Adapted from Residual Stress Measurement by X-ray Diffraction; 2003 ed., SAE International: Warrendale, PA, USA, 2003.

Figure 19 Residual stress profiles in induction-hardened 1045 steel fatigue specimens before fatigue testing and after fatigue testing for 15 million cycles at the endurance limit. Residual stress data are for bars that were initially (a) normalized and (b) quenched and tempered before induction hardening. For each residual stress profile, the best fit line and the 99% confidence interval are shown (reproduced from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

The strain (34j) at each of the sin2 j values is then measured while holding 4 and q constant (Figure 17). These values are then plotted as shown in Figure 18. The two principal stresses, s1 and s2, can be determined by graphing sin2 j versus 34j and with the knowledge of Poisson’s ratio (y) and the elastic modulus (E) for the material of interest, as shown in Figure 18 (23). Using elasticity equations, the third principal stress, s3, can also be determined. During cyclic loading, plastic deformation of the part results in relaxation of the residual stresses (27–29). For example, Figure 19 shows the residual stress profile generated in a medium-carbon steel, induction hardened from two different starting microstructures, immediately after induction processing and after fatigue cycling at the endurance limit. The two starting microstructures of the steel were normalized ferrite-pearlite and quenched and tempered martensite. Some relaxation of the residual stress was observed in the 1045 normalized condition, likely because the core microstructure undergoes some plastic deformation even at stresses as low as the endurance limit. However, the 1045 quenched and tempered condition, which is not as susceptible to plastic deformation due to its high strength, exhibited little to no relaxation of the residual stresses. It is expected that after fatigue cycling at stresses much greater than the endurance limit (low cycle regime), a significant amount of residual stress relaxation will occur.

12.16.4

Characterization of Induction-Hardened Components

Characterization of the workpiece after induction hardening is critical to determine whether the power, frequency, and scan rate settings are correct for the target induction-hardened depth. The quickest method to determine the case depth is to section the part,

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Table 2

Macro and micro etchants to determine visually the effective case depth after induction hardening

Type of etchant

Recipe

Description

Figure

Macro

10 ml nitric acid (HNO3) 90 ml ethanol 10 ml hydrochloric acid (HCl) 90 ml ethanol 2 ml nitric acid (HNO3) 98 ml ethanol 4 g picric acid (HNO3) 100 ml ethanol

Good etching response to distinguish case and core regions

(Figure 20a)

Adequate etching response between case and core

(Figure 20b)

Adequate etching response between case and core

(Figure 20c)

Good delineation of the case/core transition region

(Figure 20d)

Macro Micro Micro

593

rough grind (i.e., 320 grit SiC paper), and etch with any of the solutions listed in Table 2. The etchants attack the case martensite microstructure differently than the core, so the case depth can be measured with a Brinell scope or stereo microscope. Examples of etching with each of the solutions are shown in Figure 20 on an induction-hardened medium-carbon steel bar. For this alloy, the 10% nitric, 90% ethanol solution is the most effective at discerning the case and core. It should be noted that even though the macro etch reveals the case and core due to differences in microstructure, the most accurate method of determining the case depth is by microhardness measurements on a polished cross section. Either Vickers or Knoop microhardness measurements may be employed to measure the case depth. Determination of the case depth by hardness requires more surface preparation than the visual method; a final 1 mm diamond polish is necessary to minimize variations in measurements. The highest load possible should be used for the hardness measurements in accordance with ASTM E384-11 ‘Standard Test Method for Knoop and Vickers Hardness of Materials’ (31). The microhardness profile indicates useful information about induction-hardening processing such as the extent of transformation and excess heating of the microstructure in the case/core transition region. An ideal induction-hardened microhardness traverse is shown in the schematic in Figure 21(a). When the case is fully transformed to martensite (i.e., it was fully transformed to austenite before quenching), the hardness is constant in the case region, as shown in Figure 21(a). There is also a relatively abrupt

Figure 20 Macroetching response of an induction hardened shafts to (a) 10% nitric acid in ethanol, (b) 10% hydrochloric acid in ethanol, (c) 2% nitric acid in ethanol, and (d) 4% picric acid in ethanol.

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 21 Microhardness profiles for determination of effective case depth. An ideal hardened profile is shown in (a), an ideal hardened profile with a heat-affected zone is shown in (b), and (c) demonstrates incomplete austenitization of the parent microstructure in the case during hardening.

case/core transition region, and the hardness at the end of the transition region of the case–core interface is near the hardness of the core microstructure. Tempered martensite core microstructures are susceptible to overtempering during induction processing, which results in decreased hardness at the end of the case–core interface transition region as shown in Figure 21(b). If austenitization of the parent microstructure is not completed during induction processing, then a downward slope in hardness would be observed in the case as well as a lengthy case/core transition region (Figure 21c). The surface hardness is similar in both Figures 21(a) and 21(c), but the case hardness profile in Figure 21(a) is more desirable; thus, surface hardness measurements may not be sufficient to evaluate the effectiveness of induction heat treatments. The effective case depth by microhardness can be easily acquired after the hardness traverse is plotted similar to Figure 21. It is common in industry to define the effective case depth as the distance from the surface to the location that hardness falls below 50 HRC. However, this definition does not take into account carbon content. A more comprehensive definition of effective case depth is found in JIS G0559 “Steel – Determination of case depth after flame hardening or induction hardening” (32), which defines the effective case depth as:

l

0.23–0.33 wt. pct. C

350 HV (36 HRC)

l

0.34–0.43 wt. pct. C

400 HV (41 HRC)

l

0.44–0.53 wt. pct. C

450 HV (45 HRC)

l

0.54 wt. pct. C and greater

500 HV (49 HRC)

After microhardness measurements, the polished samples should be etched to observe the extent of microstructural transformation during hardening. For example, Figure 22 shows the microstructure from an induction-hardened part that did not fully austenitize in the case during processing. Small ferrite grains are visible in the case region, and the volume fraction of ferrite is higher as the depth increases. The parent microstructure was composed of ferrite and pearlite. The partial transformation could be due to

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Figure 22 Optical micrographs of the microstructural gradient from the surface (left) to the core (right) of an induction-hardened 1045 normalized condition with higher magnification images of the case, transition, and core regions, 4 pct. picric acid in methanol. Note the retained ferrite at the surface of the part (reproduced from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

Figure 23 Optical micrographs of the microstructural gradient from the surface (left) to the core (right) of a 4145 as-hot rolled condition with higher magnification images of the case, transition, and core regions, 4 pct. picric acid in methanol (reproduced from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

either insufficient temperature or time at temperature. The transformation kinetics could be aided through a normalizing heat treatment followed by a forced air cool to reduce the ferrite grain size in the parent microstructure. Figure 23, on the other hand, demonstrates a fully martensitic microstructure in the case, which means the case was fully austenitized during induction processing. There is also an abrupt transition from the case to the core microstructure, which is indicative of a good induction-hardening operation. The iterative steps to determine if a part has been induction hardened properly are: l

Macro etch to observe the extent of transformation. Microhardness to determine the true effective case depth. l Microstructural analysis to ensure that full transformation of the parent microstructure occurred. l

12.16.5

Strength and Fatigue of Induction-Hardened Parts

The torsional, axial, and bending load-carrying capacity of components can be increased by induction hardening. This concept is illustrated in Figures 24(a) and 24(b), which are schematic plots of applied torsional stress and effective torsional strength as a function of depth from the surface of a case-hardened bar. Microhardness values can be converted to torsional yield, tensile, or

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 24 (a) Schematic showing applied torsional stress and torsional strength for a surface-hardened bar with two different core hardness levels. (b) Schematic showing applied torsional stress and torsional strength for a surface-hardened bar with two different case depths.

fatigue strengths; a general conversion from Vickers hardness to reversed torsional fatigue strength is 1.13*HV (33). The torsional stresses associated with three levels of applied torque are plotted on the figures. In Figure 24(a), the torsional strength is plotted for two levels of core hardness. In the low-torque condition, the applied torsional stress does not exceed the torsional strength at any bar depth. In the medium-torque condition, the applied torsional stress is less than the torsional strength in the case but exceeds the torsional strength at the case–core interface for the lower core hardness condition. However, if the core hardness is increased, the applied stress does not exceed the material strength at any depth in the medium-torque condition. In the high-torque condition, the applied stress exceeds the torsional strength at the surface of the bar and near the case–core interface. Thus, the load-carrying capacity of a component is limited by the torsional strength at the surface. Similarly, Figure 24(b) shows the effects of increasing case depth for the lower core hardness condition. If the case depth is increased, the applied stress does not exceed the torsional strength at any depth in the medium-torque condition. In the high-torque condition, the torsional strength is exceeded at the bar

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

597

Figure 25 (a) Torsional yield and tensile strength as a function of effective case depth (depth to 40 HRC) in a variety of medium-carbon steels (reproduced from the Fett, G. A. Importance of Induction Hardened Case Depth in Torsional Applications. In Heat Treat. Prog. 2009, 9, 15–19). (b) Maximum elastic bending stress as a function of case depth (depth to 50 HRC) in 1541 and 4140 steel with a bar diameter of 12.7 mm.

surface for both case depths, which again shows there is an upper limit for increasing load capacity that is defined by the torsional strength at the surface. The above arguments are supported by experimental data presented in Figures 25(a) and 25(b). Figure 25(a) shows torsional strength versus effective case depth (depth to 40 HRC) in a variety of steels with carbon contents approximately between 0.4 and 0.5 wt. pct (34). Both yield and ultimate strength increase as case depth increases, but an upper limit is reached at an approximate case depth of 25% of the bar diameter. Scatter in the data is due to different core hardness levels and total case depths (depth to 20 HRC). Similarly, Figure 25(b) shows that the maximum elastic bending stress increases with increasing case depth (depth to 50 HRC) in induction-hardened 1541 and 4140 steel (18). Another objective of induction hardening is to increase the fatigue resistance of components. As described previously, compressive residual stresses arise in the case, which are beneficial for fatigue resistance. The hard martensitic case increases fatigue resistance at the surface due to the increase in local strength level. Often, the core microstructure is more susceptible to fatigue damage than the case. However, in components subjected to bending or torsion loads, higher loads are necessary to cause fatigue damage in the core region than if the same microstructure was present at the surface in a noninduction-hardened component. This latter point is illustrated in Figure 26. Figure 26(a) is a plot of fatigue data from 1040, 1541, and 4140 induction-hardened steel shafts subjected to a stress amplitude of 407 MPa in torsional fatigue (34). The figure shows that the fatigue life increases as effective case depth, defined as the depth to 40 HRC, as a percentage of the shaft diameter increases. Similarly, Figure 26(b) is a plot of bending stress amplitude versus number of cycles to failure in induction-hardened calcium-treated 4140 steel (35). Three different case depths were assessed: 22, 29, and 39% of the bar diameter. The endurance limit increases as case depth increases.

Figure 26 (a) Cycles to failure versus effective case depth as a percentage of shaft diameter for 1040, 1541, and 4140 steel shafts subjected to a stress amplitude of 407 MPa in torsion (reproduced from Fett, G. A. Importance of Induction Hardened Case Depth in Torsional Applications. In Heat Treat. Prog. 2009, 9, 15–19.). (b) Bending stress amplitude versus cycles to failure for calcium-treated 4140 steel induction hardened to 22, 29, and 39% of the bar (reproduced from the Hayne M. L. The Effect of Ferrite and Pearlite Banding on the Rotating Bending Fatigue Behavior of Induction Hardened 4040Ca4 Bar Steel. M.S. Thesis, Colorado School of Mines, Golden, CO, USA, 2010).

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

12.16.5.1 Estimation of Fatigue Crack Nucleation Location Components are typically designed to operate in the low-cycle fatigue regime where the majority of the fatigue life is spent initiating a fatigue crack. Thus, it is important to determine the regions of a component that are susceptible to fatigue crack nucleation, which is dependent on the case depth. A useful tool to determine the optimal case depth is the Woodvine diagram (36). These diagrams were first used to determine ideal carburizing depths but can also be applied to induction-hardened components (37). They can determine the most likely locations for fatigue crack nucleation in a part, and both bending and axial loads can be analyzed. The location of maximum stress, whether on a smooth specimen or at a stress concentration, should be analyzed. Similar to Figures 24(a) and 24(b), the Woodvine diagram contains the applied stress from the surface of the part to the center; the applied stress is constant for a uniaxial loading condition but varies linearly from the surface to the center for torsional or bending loads. If available, the residual stress as a function of depth should be added to the applied stress to obtain an effective applied stress (residual plus applied stress). The local fatigue strength as a function of depth is superimposed on the same plot. The local fatigue strength can be estimated using data from a microhardness traverse and a relationship such as the following: Fatigue strength ðMPaÞ ¼ 1:3695 ðHVÞ þ 48:265

[6]

where HV is Vickers hardness (33). Equation [6] is plotted along with the 95% confidence interval in Figure 27 for a variety of lowand high-alloy steels ranging from 0.25 to 0.55 wt pct. C. The locations of the part susceptible to fatigue damage are predicted by the regions on the Woodvine diagram where the effective applied stress is greater than the estimated fatigue strength (Figure 28). A ‘safe’ operating load for a given case depth or ideal case depth for a given load can be determined for a case-hardened component. The analysis should consider that the fatigue strengths are only estimated values. There are four possible outcomes from the Woodvine diagram. The first outcome (Figure 28(a)) is that there are no regions susceptible to fatigue crack nucleation. The second outcome (Figure 28(b)) is the surface of the part is susceptible to fatigue crack nucleation. In this scenario, either a lower load or a higher carbon content (which will result in greater case hardness) should be employed to increase fatigue resistance of the case. The third outcome is that both the surface and the core are susceptible to fatigue crack nucleation (Figure 28(c)). In this scenario, both higher carbon content and a deeper case depth should be considered if the applied load cannot be reduced. The final outcome is that the core is susceptible to fatigue crack nucleation (Figure 28(d)). To improve fatigue resistance for this last scenario, either a deeper case depth or higher core hardness should be considered. The latter three outcomes are for applied stresses greater than the endurance limit. Woodvine diagrams for fatigue specimens from 1045 steel, quenched and tempered before induction hardening, and 4145 steel, as-hot rolled before induction hardening, are shown in Figures 29 and 30; the fatigue samples were tested in fully reversed cantilever bending fatigue. The depths where fatigue crack nucleation was experimentally observed are also plotted on the figures. The applied stress profiles on the diagrams correspond to the stress amplitude at the experimentally determined endurance limit. The Woodvine diagrams correspond to the scenario presented in Figure 28(d). In both the diagrams, the applied stress plus residual stress at the surface of the part is zero or slightly compressive. All of the fatigue cracks nucleated in the subsurface of the fatigue

Figure 27 Relation between endurance limit and Vickers hardness for a variety of steel alloys (different alloying content and carbon content) (reproduced from the Nishijima, S. Basic Fatigue Properties of JIS Steels for Machine Structural Use; NRIM Special Report (Technical Report) No. 93–02, 2-3-12 Nakemeguro, Meguroku, Tokyo, Japan, 1993).

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

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Figure 28 Locations are susceptible to fatigue failure when the estimated fatigue strength is less than the effective stress. Based on the diagram and any stress concentrators in the part either (a) no area is susceptible, (b) the surface of the part is susceptible, (c) the surface and the core are susceptible, or (d) the core is susceptible to fatigue crack nucleation.

Figure 29 Woodvine diagram for the 1045 quenched and tempered induction-hardened condition, with the applied stress plotted at the endurance limit. Fatigue crack nucleation depths, independent of stress, are shown for all fractured samples and are separated into microstructural feature (X) and inclusion ( ) nucleated fatigue cracks (reproduced from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

600

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 30 Woodvine diagram for the 4145 as-hot rolled induction-hardened condition with the applied stress plotted at the endurance limit. Fatigue crack nucleation depths, independent of stress, are shown for all fractured samples and are separated into microstructural feature (X) and inclusion ( ) nucleated fatigue cracks (reproduce from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

specimens, which is due to the low effective applied stress at the surface. Even if the residual stress profile is not considered, the applied stress at the endurance limit does not exceed the estimated fatigue strength at any point from the case to the core. There is excellent correlation between the measured depths of fatigue crack nucleation and the regions predicted to be susceptible to fatigue crack nucleation from the Woodvine diagrams. It should also be noted that even at the endurance limit, the effective applied stress in the core exceeds the estimated fatigue strength of the material in both the 1045 quenched and tempered and 4145 as-hot rolled conditions. However, fatigue cracks may not propagate if the stress in the core is not great enough to exceed the threshold stress intensity (DKt) required to propagate a nucleated fatigue crack. However, at just 5 MPa above the endurance limit, finite fatigue lives were observed in these alloys, which indicates that the DKth required to propagate a fatigue crack was exceeded.

12.16.5.2 Fatigue Fracture Surface Location and Morphology Fatigue crack nucleation location can vary depending on the applied stress, type of loading, local hardness of the part, case depth, and resulting residual stress profile after induction hardening. The Woodvine diagram described in the previous section is useful for determining the regions of a fatigue part susceptible to fatigue crack nucleation, but many other metallurgical factors can affect fatigue life. For example, the inclusion population in the steel, the service conditions, the hydrogen content of the steel, incorrect processing parameters (during induction hardening, tempering, or both), and surface damage on the part can adversely affect fatigue life. Figure 31 shows fracture surfaces from induction-hardened bending fatigue specimens; Figure 31(a) shows a 1060 steel fracture surface, while Figure 31(b) shows a 1045 steel fracture surface. Both fracture surfaces exhibit subsurface crack nucleation with some characteristic fracture features. The region of crack initiation and fatigue crack growth is the flat, oval-shaped region at the case–core interface (arrows in Figures 31(a) and 31(b). A higher magnification image of the subsurface crack initiation region is shown in Figure 31(c). Radial marks point back to the crack initiation location. Some have attributed the flat fracture in the vicinity of crack nucleation to hydrogen embrittlement. Once the fatigue crack reaches a critical length, likely correlating to where the fracture toughness of the material is exceeded, overload fracture occurs. The overload fracture morphology is distinct between the case and core microstructures. In the brittle case microstructure, both intergranular and quasi-cleavage may be present (30,38), resulting in a flat macroscopic appearance. However, a very small shear lip region may be observed at the edge of the case region in Figure 31(c). In the core, cleavage or microvoid coalescence may be observed depending on the core microstructure (30,38). The failure is more ductile in the core and has a much rougher macroscopic appearance. Surface nucleated failures are also possible in bending fatigue (39), but compressive residual stresses at the surface and the comparatively weak core microstructure make subsurface crack nucleation more likely. Torsional fatigue fractures of induction-hardened components also exhibit distinct overload regions in the case and core microstructures. However, fatigue crack initiation appears different because of the different loading condition, and surface nucleated failures occur more readily. Figure 32 shows a schematic of a shaft subjected to torsion loading and two stress elements within the

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

601

Figure 31 Fracture surfaces of (a) a 1060 induction-hardened bar with a normalized core, (b) a 1045 induction-hardened bar with a quenched and tempered core, and (c) higher magnification image of the same 1045 bar showing the crack initiation location. (Image courtesy of D.K. Matlock).

Figure 32

Schematic of a shaft subjected to a torsional load with stress elements showing the orientation of the maximum shear and normal stresses.

shaft showing the orientations with the maximum shear and normal stresses. The maximum shear stress is transverse and longitudinal with respect to the axis of the shaft, while the maximum normal stress is oriented 45 to the axis of the shaft. Cracks nucleate either on the plane of maximum shear or maximum normal stresses. At high stress or strain amplitudes where there is significant plasticity, cracks nucleate on the plane of maximum shear stress and strain. At lower stress or strain amplitudes, cracks nucleate on the plane of maximum normal stress. Figure 33 schematically shows observed crack nucleation mechanisms in the low-, intermediate-, and high-cycle regions of a strain life plot and actual fatigue cracks in induction-hardened shafts tested at high and low shear strain amplitudes (40). The schematic shows that as the applied shear strain amplitude decreases, the crack orientation changes from predominately longitudinal (plane of maximum shear stress) to 45 (plane of maximum normal stress) from the specimen axis. As presented previously, components are designed to operate in the high-cycle fatigue regime, where crack initiation is the majority of the fatigue life. Crack initiation in induction-hardened components can occur due to plastic damage in the microstructure or around inclusions, which act as small-stress concentration regions within an alloy. Figure 34(a) shows a scanning electron microscope (SEM) image of crack nucleation due to plastic damage in the ferrite-pearlite core microstructure of a 1045 induction-hardened steel bar subjected to bending fatigue. Figure 34(b) shows an SEM image of crack nucleation due to a nonmetallic inclusion in the tempered martensite core microstructure of another 1045 induction-hardened bar subjected to bending fatigue. In softer core microstructures, such as ferrite–pearlite microstructures, fatigue crack nucleation due to plastic damage in ferrite grains is more frequent. In harder core microstructures, such as quenched and tempered martensite, the regions around inclusions are the weak links of the microstructure. Nonmetallic inclusions are present in all steels. Some examples of common nonmetallic inclusions are MnS, CaS, Al2O3, and MgO. There are many other nonmetallic inclusions present in steel, and inclusions can be composite in nature comprising one or more compound. Current steel-melting practices can produce a high-quality low-inclusion content steel, but inclusions cannot be completely removed from steel.

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Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

Figure 33 Schematic showing observed fatigue crack orientations in 1050 induction-hardened steel shafts at low-, intermediate-, and high-shear strain amplitudes. Two examples of fatigue cracking are shown below the schematic. Data replotted from the Cryderman, R.; Shamsaei, N.; Fatemi, A. Effects of Continuous Cast Section Size on Torsion Deformation and Fatigue of Induction Hardened 1050 Steel Shafts. J. Mater. Process. Technol. 2011, 211, 66–77.

Figure 34 Examples of fatigue crack nucleation at the case–core interface in induction-hardened bars subjected to bending fatigue (a) microstructural features in a ferrite/pearlite 1045 induction-hardened steel where plastic damage likely contributed to crack nucleation and (b) an inclusion at the crack nucleation location in a 1045 quenched and tempered steel (reproduced from the Nissan A. B. The Effect of Microstructure and Induction Processing on Fatigue Performance and Crack Initiation of Induction Hardened Bar Steel. Ph.D. Dissertation, Colorado School of Mines, Golden, CO, USA, 2012).

12.16.6 l l l l l l

Industrial Standards Relevant to Induction Hardening

International Standard (ISO) 3754 “Steel – Determination of Effective Depth of Hardening after flame or induction hardening” Japanese Industrial Standard (JIS) B 6912:2002 “Process of Induction Hardening and Tempering of Iron and Steel” Japanese Industrial Standard (JIS) G 0559:2008 “Steel – Determination of Case Depth after Flame Hardening or Induction Hardening” Society of Automotive Engineers (SAE) ARP4715 “Induction Hardening of Steel Components” ASTM Standard A255-10, “Standard Test Methods for Determining Hardenability of Steel” ASTM Standard E384-11, “Standard Test Methods for Knoop and Vickers Hardness of Materials”

Microstructures and Mechanical Performance of Induction-Hardened Medium-Carbon Steels

12.16.7

603

Summary and Conclusions

Induction processing is an effective method to produce case-hardened components in low- and medium- carbon steels with various microstructures. The induction processing parameters can be carefully tuned to achieve desired case depths. The induction-hardening process and alloy selection have interrelated effects on the resulting mechanical behavior, and the important parameters to consider are case depth, residual stress distribution, case and core microstructure hardness, and other microstructural features such as inclusions. There are many opportunities for further research and development in induction-hardened steels. The development of the induction-hardened case microstructure as a function of induction process parameters and alloy should be characterized, especially with more widely available high-resolution characterization equipment such as advanced transmission electron microscopes and atom probe tomography. The role of microalloying in microstructural development and its effects on mechanical performance should also be assessed. For example, micoralloying may increase the fatigue resistance of the core microstructure and thus improve mechanical performance. Several of the fatigue research studies conducted have utilized bending fatigue tests, and crack nucleation often occurs at the case–core interface; thus, fatigue performance is controlled by the core microstructure. In contrast, torsional loads are applied to many induction-hardened part applications, and fatigue crack nucleation occurs at the surface of the part; then, the case microstructure plays a larger role. As discussed in this chapter, torsional fatigue studies have been performed, but further work is necessary to assess the effects of microstructure in the case and core on torsional fatigue performance.

Acknowledgments The authors are grateful for the support, in writing this chapter, of the Advanced Steel Products and Processing Research Center, an industry–university collaborative research center at the Colorado School of Mines.

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