defects and distortion in heat-treated

Defects and Distortion in Heat-Treated Parts Anil Kumar Sinha, Bohn Piston Division

MOST OF THE PROBLEMS in heat- treated parts are attributed to faulty heat- treatment practices (such as overheating and burning, and nonuniform heating and quenching), deficiency in the grade of steels used, part defect, improper grinding, and/or poor part design. This article discusses overheating and burning, residual stresses, quench cracking, and distortion in some detail and offers some suggestions to combat them.

Most of these conditions result in a charac- teristic appearance of the treated parts that can be easily recognized by simple inspection. Some of these factors do not produce any distinguishing features in the semifin- ished or finished part. In particular, some of the visual evidence does not recognize the presence of overheating and burning and the development of residual stresses leading to distortion, quench cracking, and eventual failure of the heat-treated parts; metallurgical laboratory examination is needed to establish these problems that contribute significantly to the service performance of the part. Tool designers must also be aware of the problems and difficulties in manufacture, heat treatment, and use.

Overheatin 8 and Burning of Low-Alloy Steels

When low-alloy steels are preheated to high temperature (usually > 1200 C, or 2200 F), prior to hot mechanical working (such as forging) for a long period, a deterioration in the room-temperature mechanical properties (particularly tensile ductility and impact strength or toughness) can be obtained after the steel has been given a final heat treatment (comprising reaustenitizing, quenching, and tempering) (Ref 1-3). Linked with the im- paired mechanical properties is the appearance of intergranular matte facets on the normal ductile fracture surface of an impact specimen. This phenomenon is known as overheating and has been a matter of con- cern, especially in the case of steel forgings. Overheating has also been noticed in steel castings (due to variation in pouring temper-

ature and effectiveness of the proprietary grain inoculants applied to the mold surface), in heavily ground parts, and in affected zones of welds (Ref 4). The usual practice is to reject the overheated products as being un- suitable for service.

It has now been established that overheating is essentially a reversible process caused by the solution of MnS particles in austenite during heating or reheating at high temperatures; the amount increases with temperature, and its subsequent reprecipi- tation during cooling occurs at intermediate rates as very fine ( -0 .5 i~m) arrays of a-MnS particles on the austenite grain boundaries. On subsequent heat treatment the intergranular network of sulfides may provide a preferential, lower-energy fracture path in contrast to a normal transgranular fracture path. As a result, when impact loaded, a ductile intergranular fracture develops due to decohesion of the MnS/matrix interface and progress of microvoid coales- cence. Figures 1 (a) and (b) show the usual appearance of the fracture surface at different magnifications (Ref 1).

When the low-alloy steel is preheated prior to hot working at too high a temperature (normally > 1400 C, or 2550 F), local melting occurs at the austenite grain boundaries as a result of the segregation of phosphorus, sulfur, and carbon (Ref 5). During cooling, initially dendritic sulfides (proba- bly type II-MnS) form within the phosphorus-rich austenite grain boundary, which then transforms to ferrite. This results in excessively weak boundaries. Subsequent heat treatment provides a very poor impact strength and almost completely intergranular fracture surface after impact failure. This phenomenon is termed burning. Burn- ing thus occurs at a higher temperature than overheating. If this occurs during forging, the forging will often break during cooling o r subsequent heat treatment (Ref 4).

Detection of Overheating There are two basic methods for the

determination of the occurrence of over-

166,6 ~rn

I I 12.5 p, rn

Fracture surface of an impact loaded speci- Fig 1 men. (a) Appearance of intergranular fracture of 4.25Ni-Cr-Mo steel containing 0.34% Mn and 0.008% S, in fully heat-treated condit ion but after cooling from 1400 C (2550 F) at 10 C/min (20 F/rain). (b) Same specimen as in (a) but at higher magnifica- t ion, showing ductile dimples nucleated by MnS particles precipitated at austenite grain boundaries. Courtesy of The Institute of Metals

heating, namely, fracture testing and metallography (or etch testing). Overheating may also be detected by a decrease in mechani-

ASM Handbook, Volume 4: Heat Treating ASM Handbook Committee, p 601-619

Copyright 1991 ASM International All rights reserved.

www.asminternational.org

602 / Process and Quality Control Considerations

Table 1 Etching characteristics of overheated and burned steels Reagent Method Action on overheated steel Action on burned steel

2.5% nitric acid in ethyl Swab surface for 30 s May produce grain contrast, but White boundaries outlining alcohol not indicative of overheating preexisting austenite grains

Saturated aqueous solution White boundaries outlining Black boundaries outlining of ammonium nitrate preexisting grains preexisting austenite grains

Aqueous 10% nitric acid + 10% sulfuric acid

85% orthophosphoric acid (Fine's reagent)

Oberhoffer's reagent

Black boundaries outlining preexisting austenite grains

Does not differentiate between overheated and nonoverheated steel

Does not differentiate between overheated and nonoverheated steel

Source: Ref 13

Electrolytic, specimen anode, current density 1.0 A cm -2 (6.5 A in. -2)

Etch for 30 s, swab surface; repeat three times, then repolish lightly

Electrolytic, specimen anode, current density 0.15 A cm -2 (1.0 A in.-2), etching time 15 min

Swab surface for 30 s

White boundaries outlining preexisting austenite grains

Attacks inclusions at grain boundaries

Shows phosphorus segregation at grain boundaries

cal properties. But such changes are not very marked unless overheating temperature is high or overheating is too prolonged or severe; in some instances the mechanical properties do not change, even after the observation of extensive faceting. Usually the two methods mentioned above should be used in conjunction with some measure of toughness by impact or other testing in order to get a clear understanding of the degree and severity of overheating (Ref 2).

Fracture Testing. The direction of fracture testing is important in steels manufactured by conventional methods. It has been observed by some workers (Ref 6) that the longitudinal fracture test specimens parallel to the rolling direction do not exhibit face- ring until the corresponding transverse frac- tures display extensive faceting. However, the testing direction in electroslag-refined (ESR) steels has been found to be insignif- icant (Ref 7).

The scanning electron microscope is considered to be the best and most convenient tool to detect the facets on the overheated fracture surfaces. These facets are charac- terized by small, well-defined, ductile dimples; each dimple is usually nucleated, presumably by fine arrays of inclusion particles: a-MnS particles (Fig 1) in Mn- bearing steels (Ref 8, 9) or chromium sulfides in Mn-free steels (Ref 10, 11).

It is now well recognized that the fracture test specimen should always be tested in the toughest possible state (for example, quenched and highly tempered [in the range 600 to 650 C, or I 110 to 1200 F] steels after high-temperature austenitization) because this condition is most prone to overheating effects. Baker and Johnson (Ref 5) have suggested that an increased proportion of facets in the fracture specimens with increasing tempering temperature is attributed to the corresponding increase of the plastic zone size. In this case a slight amount of weakening will be sufficient to impart faceting because the grain boundary strength becomes lower (Ref 2). It should be noted that the existence of facets in the fractured specimens is not always associat-

ed with a lowering of impact strength (Ref 12).

Metallography (or Etch Testing). The most widely used etchant technique uses Austin's reagent (aqueous solution of 10% nitric and 10% sulfuric acids), ammonium persulfate, molten zinc chloride, saturated solution of picric acid at 60 C (140 F), and an electrolytic etch based on saturated aqueous ammonium nitrate. Table 1 shows the etching characteristics of overheated and burned steels (Ref 13). The etchant procedure with Austin's etchant is as fol- lows: The sectioned specimen is etched for 30 s in the etchant, removed, washed off, and repeated three times. If the steel has been overheated, the original austenite grain boundaries will be preferentially at- tacked, and a black network of etch pits will be observed under the microscope (Ref 14). According to Preece and Nutting (Ref 13), the best results are obtained when ammonium nitrate etch is applied on the sectioned steel specimen in the fully heat-treated condition where this etchant preferentially attacks the matrix (original austenite grains), leaving the grain boundary unaffected (which appears as a white network). Bodimeade (Ref 15) concluded that all these etchants did not cope with mildly overheated low-sulfur steels. Table 2 is a summary of the results of potentiostatic etching techniques carried out by McLeod (Ref 12) using nitric-sulfuric, saturated aqueous picric acid (at 60 C, or 140 F), and ammonium nitrate etchants. He considered that when the suitable etching conditions were established, the potentiostatic etching method rendered more reliable and reproducible results as compared with the conventional etching techniques. However, the same problem with mildly overheated low-sulfur steels still persisted. Hence, the use of etch tests for low-sulfur low-alloy steels is not recommended for the detection of mild overheating.

Detection and Effects of Burning Burning is not commonly encountered.

The two etchants (namely, nitric-sulfuric

acid and ammonium nitrate solution) used for overheating can be successfully employed for detecting burning. When applied to burned steels, these etchants react in a manner opposite to that of overheated steels. Preece and Nutting (Ref 13) found ammonium nitrate solution to be the ideal reagent to detect this phenomenon. Other reagents are Stead's and Oberhoffer's reagents, which may also be used to check the burning effect. However, these etchants are unable to differentiate between overheated and nonoverheated steels.

Factors Affecting Overheating The occurrence and severity of overheat-

ing depend principally on important factors, notably steel composition, temperature, cooling rate, and method of manufacture.

Composition. Sulfur is the constituent that greatly influences overheating. For steels with less than 0.002 wt% sulfur, overheating does not occur; this is because of the very low volume fraction of sulfides formed. However, the commercial production of such very-low-sulfur steels (for example, ESR steels) is expensive. Above this level of sulfur, the overheating onset temperature rises with the increasing amount of sulfur. It has now been explained that steels with low sulfur content (0.01 to 0.02%) are more prone to this defect than those with high sulfur content (>0.3%) because the transgranular strength is high, and therefore a small amount of grain-boundary sulfide precipitation is enough to induce intergranular failure (Ref 16). The phosphorus content has been regarded with the most con- cern in connection with burning. At constant phosphorus level, there is an increase in the overheating temperature with the increase of sulfur content, whereas the burning onset temperature decreases. Burn- ing temperature is reduced with the increase in phosphorus content. At low sulfur contents, a wide gap between overheating and burning temperatures exists. For example, in the case of vacuum remelted steels, the temperature gap between the onset of overheating and burning is -300 to 400 C ( -570

Defects and Distortion in Heat-Treated Parts / 603

Table 2 Summary of potentiostatic etching experiments Best etching conditions

Anodic loop Solution voltage, mV Observed effect Voltage, mV Observed effect Comments

Saturated aqueous -400 Slight general 2200 (for 2 Classic white ammonium etching min) boundaries on a nitrate dark background

Aqueous 10% nitric acid + 10% sulfuric acid

Saturated aqueous picfic acid at 60C (140 F)

Source: Ref 12

200 Vigorous None dissolution of specimen; formation of flaky black film

-250 Milder attack; About -250 large black (for 30 s) pits in mildly etched matrix

100 No real, None positive indication of overheating

Discontinuous array of grain-boundary pits and some random pits within grains

Operates best in the transpassive region at >+1500 mV; time at any potential is important

Underetching: random array of black pits

Overetching: uniform black surface film

Most aggressive etchant of the three examined

Polish lightly after etching to eliminate matrix etching effects

Anodic loop very weak, necessitating long etching times because current density is very low; Teepol additions gave no improvement

to 750 F) and there is a remote possibility of burning occurring within the forging range, unless the overheating is severe (Ref 2). However , at high sulfur content the gap becomes narrow.

Temperature. To avoid overheating, care must be exercised in choosing a correct heating temperature so that uneven heating, flame impingement, and so forth, do not occur (Ref 3).

Cooling Rates. The cooling rate through the overheating range affects the size and dispersion of intergranular et-MnS particles. The intermediate cooling rate generally employed, 10 to 200 C/min (20 to 360 F/min), gives rise to maximum faceting as well as to the greatest loss in impact strength. How- ever, slow and rapid cooling rates will sup- press overheating. At very slow cooling rates, the sulfide particles become large, small in number, and more widely dispersed, and they have no more deleterious effects than the other inclusions already present. At rapid rates, the sulfide inclusions are too fine to produce any damaging effect (Ref 17).

Methods of Manufacture. Electroslag- remelted steels are less susceptible than vacuum-remelted steels, presumably due to the difference in oxygen level. Similarly, nickel steels are more prone to overheating. Vacuum-remelted steels have a lower overheating temperature than some comparable air-melted steels.

Prevention of Overheating and Burning

For preventing overheating of steels, a properly selected temperature should lie

between a temperature low enough for the metal to be safe and high enough to be sufficiently plastic. The better the temperature control, the better the compromise.

Severe overheating can be reduced to mild overheating by soaking the steel at 1200 C (2200 F); with care, it may be removed completely. Hot working through the overheating range to a low finish temperature is also reported to remove the effects of overheating.

The alloying additions with a greater sulfide-forming tendency, such as calcium, zir- conium, cerium ( -0 .3% of the melt), or mixed rare earth metals (in the form of misch metal containing 52% Ce, 25% La, and 12% Nd), have been shown to increase significantly both the overheating temperature and mechanical properties of the steel (for example, ductility and toughness). Pro- vided that a high Ce/S ratio (>2) existed, a complete change in sulfide morphology occurred in low-alloy steels where the elon- gated MnS inclusion occurring in the un- treated steel was totally replaced by small globular type-I rare earth sulfides and ox- ysulfides of high thermal stability even after austenitizing at 1400 C (2550 F) (Ref 2). This treatment does not show intergranular faceting. Burning can also be avoided in the same way by treating with calcium, zirconi- um, cerium, or mixed rare earth addition to form refractory, less-soluble sulfides.

Control of Cooling Rates. Control of cooling rates is not a practical method for large forgings because extremely slow cooling is prohibitively time consuming and causes excessive scaling and decarburization, and rapid quenching from high temperatures

produces cracking and distortion of the parts (Ref 2).

Reclamation of Overheated Steel Severely overheated steels can often be

completely restored by any of the following heat treatments:

Repeated normalizing (as many as six) starting at temperatures 50 to 100 C (90 to 180 F) higher than usual, followed by a standard normalizing treatment (Ref 2)

Repeated oil-hardening and tempering treatments after prolonged soaking at 950 to 1150 C (1740 to 2100 F) in carburizing atmosphere. Rehardening more than three times is not advisable

Soaking at 900 to 1150 C (1650 to 2100 F) for several hours. This causes growth of MnS particles by the Ostwald ripening process and results in an excessive scale formation and a loss of dimensional accu- racy of the forgings

Residual Stresses

Heat treatment often causes stress- and strain-related problems such as residual stress, quench cracks, and deformation and/ or distortion. The residual stress may be defined as the self-equilibrating internal or locked-in stress remaining within a body with no applied (external) force, external con- straint, or temperature gradient (Ref 18, 19). There are two types of residual stresses:

Macro- or long-range residual s tress is a first-order stress that represents an average of body stresses over all the phases in polyphase materials. Macroresidual stresses act over large regions as compared to the grain size of the material. Traditionally, engineers consider only this type O f residual stress when design- ing mechanical parts

Microresidual stress, also t e rmed tesse- lated stress or short-range stress is a second-order or texture stress, which is associated with lattice defects (such as vacancies, dislocations, and pile-up of dislocations) and fine precipitates (for example, martensite) (Ref 20-22). Microre- sidual is the average stress across one grain or part of the grain of the material. This information is indispensable in studying the essential behavior of material deformation

These two types of residual stresses may also be classified further as a tensile or compressive stress located near the surface or in the body of a material. This section focuses on the effects, development , control, and measurement of long-range residual stresses.

Effects of Residual Stress The major effects of residual stress in-

clude dimensional changes and resistance to


Surface residual stress (root of notch), ks -200 -160 -120 -80 -40 0 40

1100 8645 notch cold rolled I 8645 notch warm rolled 160 0.25 notch radius/~- I 0.25 notch radius

~i~ J / 18645 I 1045 j I - " ~ . ~ / I shot peened I untempered a~-,,,,L// j 1 4 B 3 5

825 I 1045- ~ , , " t e m p e r ~ - - 120 ~. tempered . . . . .

te ,~ ,e red \~~ , \ temperecl .~ ._E I 8630-N ~ \ .E_ -~ = 550 I I t empered~ 8630 80 N ~

Specimen X ~ / o i l quenched

6.75 ~ ~ ' ~ 1 8660 oil uenched "' i 8645 - - " ~ ~ . / 275 - I I tempered [ ~ ' ~ - ~ _ ~ 40

L_ 1.750 in. L1.550 in. diam ~ 8645 oil quenched 60 V-notch 1

diam 0.025 root radiu Compression ~--~-Tension 0 i i I I 0

-1375 -1100 -825 -550 -275 0 275 Surface residual stress (root of notch), MPa

Effect of surface residual stress on the endurance limit of selected steel. All samples were water Fig 2 quenched except as shown, and all specimen dimensions are given in inches. Source: Ref 23, 24

crack initiation. Dimensional changes occur when the residual stress (or a portion of it) in a body is eliminated. In terms of crack initiation, residual stresses can be either beneficial or detrimental, depending on whether the stress is tensile or compressive.

Compressive Residual Stress. Because residual stresses are algebraically summed with applied stresses, residual compressive stresses in the surface layers are generally helpful because the built-in compressive stresses can reduce the effects of imposed tensile stresses that may produce cracking or failure. Compressive stresses therefore contribute to the improvement of fatigue strength and resistance to stress-corrosion cracking in a part and an increase in the bending strength of brittle ceramics and glass (Ref 22).

Figure 2 shows that the endurance limit fatigue strength of selected steels increases

with the surface residual compressive stress developed by specific heat treatment and surface processing. It is also apparent that, in the presence of high compressive stress, a poor microstructure in steel samples has a small influence on good endurance limit fatigue strength (Ref 23-25). These fatigue improvements are of great significance in components, particularly where stress raisers, such as notches, keyways, oil holes, and so forth, are highly desirable in the design of components (for example, crank- shafts, half-shafts, and so on) (Ref 26). Many fabrication methods have been developed to exploit this phenomenon. Pre- stressed parts (including shrink-fits, pre- stressed concrete, interference fits, bolted parts, coined holes, wire-wound concrete pipe), mechanical surface working processes (such as shot peening, surface roil- ing, lapping, and so on) of hardened ferrous

Table 3 Summary of compressive and tensile residual stresses at the surface of the parts created by the common manufacturing processes Compression at the surface Tension at the surface

Surface working: shot peening, surface rolling, lapping, and so on

Rod or wire drawing with shallow penetration(a) Rolling with shallow penetration(a) Swaging with shallow penetration(a) Tube sinking of the inner surface Coining around holes Plastic bending of the stretched side Grinding under gentle conditions Hammer peening Quenching without phase transformation Direct-hardening steel (not through-hardened) Case-hardening steel Induction and flame hardening Prestressing Ion exchange

Rod or wire drawing with deep penetration Rolling with deep penetration Swaging with deep penetration Tube sinking of the outer surface Plastic bending of the shortened side Grinding: normal practice and abusive conditions Direct-hardening steel (through-hardened)(b) Decarburization of steel surface Weldment (last portion to reach room temperature) Machining: turning, milling Built-up surface of shaft Electrical discharge machining Flame cutting

(a) Shallow penetration refers to ~


Table 4 Changes in volume during the transformation of austenite into different phases

Change in volume, %, as a function of carbon

Transformation content ( % C)

Spheroidized pearlite - 4 . 6 4 + 2.21 (% C) ---, austenite

Austenite ~ 4.64 - 0.53 x (% C) martensite

Spheroidized pearlite 1.68 x (% C) martensite

Austenite ~ lower 4.64 - 1.43 (% C) bainite

Spheroidized pearlite 0.78 x (% C) lower bainite

Austenite ~ upper 4.64 - 2.21 x (% C) bainite

Spheroidized pearlite 0 upper bainite

Source: Ref 4

the transformation of austenite into martensite or other transformation products (Ref 27). Table 4 lists the changes in volume during the transformation of austenite into different structural constituents (Ref 28).

Thermal Contraction. The relation between the thermal stress ~th during cooling and the corresponding temperature gradient in the component is given by:

tYth = E - AT" ct (Eq

where E is the modulus of elasticity, and tx is the thermal coefficient of expansion of the material. It is thus apparent that thermal stresses are greatest for materials with high elastic modulus and coefficient of thermal expansion. Temperature gradient is also a function of thermal conductivity. Hence, it is quite unlikely to develop high-tempera-

1000 ? c w

~ 500 ~. u

~- 0 1

Water quenched 100 mm (4 in.)

specimen

1700 ou-

1100 ~

600 E

100 ~- 10 103 Time, s

_

e~ e~

E E o o

Deve lopmen t of thermal and residual stresses F i g 3 in the longitudinal direct ion in a 100 mm (4 in.) d iameter steel bar on wate r q u e n c h i n g f rom the austenitizing t empe ra tu r e , 850 C (1560 F). Transforma- tion stresses are not taken into cons idera t ion . Source: Ref 30

Table 5 Relevant physical properties in the development of thermal stresses Coefficient of

Modulus of elasticity expansion Thermal conductivity

Metal GPa psi x 10 6 1 0 - 6 / K 1 0 - 6 p F W m -1 k - l Btu in./ft 2 h F

Pure iron (ferrite) 206 30 12 7 80 555 Typical austenitic steel 200 29 18 10 15 100 Aluminum 71 10 23 13 201 1400 Copper 117 17 17 9 385 2670 Titanium 125 18 9 5 23 160

Source: Ref 29

ture gradients in good thermal conductors (for example, copper and aluminum), but it is much more likely in steel and titanium (Ref 29). Another term involving thermal conductivity, called thermal diffusivity (Dth), is sometimes used in context with temperature gradient. It is defined a s D t h = k/pc, where k is the thermal conductivity, p is the density, and c is the specific heat. It is clear that low Oth (or k) promotes large temperature gradient or thermal contraction. It should be emphasized that large size of the part and high heating or cooling rates (severity) of quenching medium also aug- ment temperature gradients leading to large thermal contraction.

Table 5 lists some of the relevant material properties that affect thermal and residual stresses (Ref 29).

Residual Stress Pattern Due to Thermal 1)Contraction. Residual stress is developed

during quenching of a hot solid part that involves thermal volume changes without solid-state phase transformation. This situation also exists when a steel part is cooled from a tempering temperature below the A t. Figure 3 shows the development of longitudinal thermal and residual stresses in a 100 mm (4 in.) diam steel bar on water quenching from the austenitizing temperature, 850 C (1560 F) (Ref 30). At the start of cooling, the surface temperature S falls drastically as compared to the center temperature C (top left sketch of Fig 3). At time w, the temperature difference between the surface and core is at a maximum of about 550 C (1020 F), corresponding to a thermal stress of 1200 MPa (80 tons/in. E) due to linear differential contraction of about 0.6%, if relaxation does not take place. Under these conditions, tensile stresses are developed in the case with a maximum value of a (lower diagram), corresponding to time w in the upper diagram, and the core will contract, producing compressive stresses with a maximum of b. The combined effect of tensile and compressive stresses on the surface and core, respectively, will result in residual stresses as indicated by curve C, where a complete neutralization of stress will occur at some lower temperature u. Further decrease in temperature, therefore, produces longitudinal, compressive residual stresses at the surface and the tensile stresses at the core, as shown in the lower right-hand diagram of Fig 3. Figure 4(a) is a schematic

illustration of the distribution of residual stress over the diameter of a quenched bar due solely to thermal contraction in the longitudinal, tangential, and radial directions (Ref 19).

The maximum residual stress attained on quenching increases as the quenching temperature and quenching power of the cool- ant are increased. Tempered glass is made by utilizing quenching techniques in which glass is heated uniformly to the annealing temperature and then surface cooled rapidly by cold air blasts. This produces compressive surface stresses to counteract any tensile bending stress, if developed during loading of the glass, thereby increasing its load-carrying capacity (Ref 31).

Residual Stress Pattern Due to Thermal and Transformational Volume Changes (Ref 32). During quench hardening of a steel (or other hardenable alloy) part, hard martensite forms at the surface layers, associated with the volume expansion, whereas the remainder of the part is still hot and ductile austenite. Later, the remainder austenite transforms to martensite, but its volumetric expansion is restricted by the hardened surface layer. This restraint causes the central portion to be under compression with the outer surface under tension. Figure 4(c) illustrates the residual stress distribution over the diameter of a quenched bar showing volume expansion associated with phase transformation in the longitudinal, tangential, and radial directions (Ref 19). At the same time during the final cooling of the interior, its contraction is hindered by the hardened surface layers. This restraint in contraction produces tensile stresses in the interior and compressive stresses at the outer surface. However, the situation as shown in Fig 4(c) prevails, provided that the net volumetric expansion in the interior, after the surface has hardened, is larger than the remaining thermal contraction. In some particular conditions, these volumetric changes can produce sufficiently large residual stresses that can cause plastic deformation on cooling, leading to warping or distortion of the steel part. While plastic deformation appears to reduce the severity of quenching stresses, in most severe quenching the quenching stresses are so high that they do not get sufficiently re- leased by plastic deformation. Consequent- ly, the large residual stress remaining may


+~

Longitudinal

I Tangential

I

I Radial

(a) (b)

.~_ T j

L = l ong i t ud ina l T = tangen t ia l R = radial

I

I i

Longitudinal

I I

i

I Tang?ntial

I

t-'--->' Rad al

(el

Schematic illustration of the distribution of residual stress over the diameter of a quenched bar in the F ig 4 longitudinal, tangential, and radial directions due to (a) thermal contraction and (c) both thermal and transformational volume changes. (b) Schematic illustration of orientation of directions. Source: Ref 19

reach or even exceed fracture stress of steel. This localized rupture or fracture is called quench cracking (Ref 32, 33).

It should be emphasized again that for a given grade of steel, both large size of the part and higher quenching speed contribute to the larger value of thermal contraction, as compared to the volumetric expansion, of martensite. In contrast , when the parts are thin and the quenching rate is not high, thermal contraction of the part subsequent to the hardening of the surface will be smaller than the volumetric expansion of martensite. Similarly, for a given quenching rate, the temperature gradients decrease with decreasing section thickness, and consequently the thermal component of the residual stress is also decreased (Ref 24).

Figure 5(a) shows the continuous cooling transformation diagram of DIN 22CrMo44 low-alloy steel exhibiting austenitic decomposition with the superimposed cooling curves of the surface and center in round bars of varying dimensions. If the large- diameter (100 mm, or 4 in.) bar is water quenched (that is, for slack quenching), martensitic transformation occurs at the surface, and pearlitic + bainitic transformations occur at the center, resulting in a residual stress pattern (top of Fig 5) similar to that due solely to thermal stress (Fig 4a).

During the rapid quenching of the medium- size (30 mm, or 1.2 in.) bar diameter, the start of bainite transformation at the center coincides approximately with the transformation of martensite on the surface. This results in compressive stresses at both the surface and center, with tensile stresses in the intermediate region (middle of Fig 5). When the smaller-diameter (10 mm, or 0.4 in.) bar is drastically quenched (for example, in brine), the entire bar transforms to martensite. This is associated with very little temperature variation between the surface and the center of the part. In this situation, tensile residual stress is developed at the surface and compressive stress at the center of the bar (bottom, Fig 5) (Ref 34, 35).

Although the shallower hardening steels exhibit higher surface compressive stresses, deep hardening steels may develop moder- ately high surface compressive stresses with severe water quenching. When these deep hardening steels are through-hardened in a less efficient quenchant, they may exhibit surface tensile stresses (Ref 24, 31). Rose has pointed out the importance of transformations of core and surface before and after the stress reversal. According to him the tensile surface residual stress occurs when the core transforms after, and the

surface transforms before, the stress reversal (Fig 4c and bot tom of Fig 5), whereas compressive surface residual stress takes place when the core transforms before, and the surface transforms after, the stress reversal (top of Fig 5). His analysis is capable of explaining complex stress patterns for various combinations of part sizes, quenching rate, and steel hardenability (Ref 21). However , the residual stress pattern in the hardened steels can be modified either with different transformation characteristics or during the tempering and finish-machining (after hardening) operations.

Residual Stress Pattern after Surface Hard- ening. In general, thermochemical and thermal surface-hardening treatments produce beneficial compressive residual stresses at the surface.

Carburized and Quenched Steels. When low-carbon steels are carburiZed and quenched, first the core transforms at high temperature (600 to 700 C, or l l00 to 1300 F) to ferrite and pearlite with the attendant relaxation of any transformation stresses. Later , the high-carbon case transforms to martensite at much lower temperature (less than 300 C, or 570 F), accompanied by volume expansion and under conditions of no (or minimum) stress relaxation. As a result, residual compressive stress is developed in the case with a maximum at the surface.

Large differences in carbon level between the case and the core determine the se- quence of phase transformation on cooling after carburizing and the resultant development of compressive residual stress in the case. Likewise, compressive residual stress in the case increases as the core carbon content decreases. Increasing case depth reduces the contribution from the low-carbon core in the development of compressive stress in the case, thereby adversely affecting the fatigue propert ies (Ref 36).

In actual practice, a maximum compressive stress develops at some distance away from the surface (Fig 6 and 7). This effect occurs because of the presence of retained austenite, the extent of which depends on steel composition, carbon content of the case, quenching temperature, and severity of quench. According to Koistinen (Ref 38) and Salonen (Ref 39) the peak compressive stress takes place at 50 to 60% of the total case depth corresponding to about 0.5 to 0.6% carbon level, which produces a low retained austenite content and martensite hardness around the maximum. Another factor that might influence this compressive residual stress profile is that the martensite formed in the lower-carbon regions of the case is of the lath type, which also affects the retained austenite content (Ref 20). The reversal sign of residual stress takes place at or near the case/core interface. Later, when Kois t inen 's theory was applied to the measured data, it appeared that the position of


1000 1830

800 , ~ 1470

600 ~ ~ 1110

400 \ ~ . _ 750

200 " ~ ~ 390 Surface Center

0 1000 1830

800 ~ ~ 1470

400 - 750 X E E

200 390 er

0 Surfacel

1000 1830

800 ( 1470

600 ~ . . _ . . . . 1110 /

400 " ' ~-~,........._. 750

nter 0

1 10 100 103 Time, s

(a) (b)

+20

-20

Distribution of residual stresses

+3

-3

m -40 -6

~ Center Surface ~ ~ +20 +3 "~

"o "o .- "~ o o

30 mm -20 diam -3

Center Surface +20 [ +3

0 0

I ~ / 0 mm -20 ~ diam -3

Center Surface + = Tensile stresses - = Compressive stresses

(a) Continuous cooling transformation diagrams of DIN 22CrMo44 steel showing austenitic decompo- Fig 5 sition with the superimposed cooling curves of the surface and center during water quenching of round bars of varying dimensions. (b) The corresponding residual stress pattern developed because of thermal and transformational volume changes. Source: Ref 34, 35

maximum compressive stress depends on severity of quenching, total case depth, steel hardenability, and so forth (Ref 21, 40). Figure 7 shows the details of generation of axial stress distribution of a carburized gear (made from deeper hardening steel) during quenching. In the early stages, the contour lines of equal stress were largely unaffected by the surface profile. Later a zone of high compressive stress distribution occurred in the central portion of the teeth, which remained until the end of the quench (Ref 37).

In nitriding, like carburizing, a compressive residual stress is set up in the surface layers. High-temperature nitriding produces a little relaxation of stresses, whereas low- temperature nitriding imparts a maximum residual stress. In nitrocarburizing, improvement in residual surface compressive stress and fatigue strength depends on the hardness and depth of diffusion zone. These properties, in turn, decrease with increasing carbon and alloy content (that is, increased hardenability). During quenching, after ni-

trocarburizing, a (macro-) compressive residual stress is produced in the compound layer and gamma prime phase (Ref 41). When nitrocarburized parts are rapidly quenched, the above properties are further enhanced (Ref 42).

In borided steel processed at 900 C (1650 F), a high compressive residual stress is developed at the surface layers (Fig 8), which consists of FeB and Fe2B phases (Ref 43); this is attributed to the lower thermal expansion coefficient and the larger specific volume in a borided layer compared to that in a ferrite matrix (Ref 18, 43).

In an induction-hardened steel part, a compressive surface residual stress is produced when wear-resistant hard martensite (with slightly lower density) is formed on the surface of a section concurrently with volume expansion while nonhardened core remains essentially unchanged (Fig 9) (Ref 44, 45). The magnitude of the compressive stress, which is affected by both thermal contraction and martensite formation, may be a considerable fraction of the yield

1.0

c" o 0.5

8

-o , c.-~_

~ o

m

.9_ tt-

Tensile

Compressive

Distance from the surface

Relationship between carbon content, re- Fig 6 tained austenite, and residual stress pattern. It shows the development of peak compressive stress some distance away from the surface. Source: Ref 20

strength, which permits the application of significantly higher stresses than could normally be possible in fatigue loading. As in the carburizing practice, the surface compressive residual stresses are usually found to increase, with depth below the surface (Ref 45) (Fig 9, Ref 44). A fairly sharp transition to a tensile state takes place near the hardness drop-off between the case and unhardened surrounding material. With an increase in distance from the steep transition, the tensile condition gradually fades away toward zero stress (Ref 44). In induction hardening, an increase in hardenability changes the depth at which transition from compressive to tensile stress occurs. The increase in the rate of heating produces an increase in the maximum compressive and tensile residual stresses without affecting the mode of stress distribution (Ref 46).

Residual Stress in Other Processing Steps. As welding progresses, the temperature distribution in the weldment becomes nonuniform and varying as a result of localized heating of the weldment by the welding heat source. During the welding cycle, comprising heating and cooling, complex strains develop in the weld metal and adjacent areas. As a result, appreciable residual stresses remain after the completion of welding. Since the weld metal and heat- affected zone contract on cooling (Fig 10a), they are restrained by the cool adjacent part. This produces tensile residual stress in


- 900

Gz = 200MPa -600

100 ~ -300

( ~100

t= 3 s 30 s

300

3 O O

Carburized SNC815

0300 -600 60~ j

l i 300 60 s

-900 -600

0 500

Distance from surface, in. 0.002 0.004 0.006

0 0 z~ ~ 0 #_

- 5 0 0

-1000 ~ ~ o '~ -1500

300 -2000 0 0 0

-2500 600 0 0.05

Fig 7 Axial stress distribution (given in MPa) in carburized gear during quenching process. Source: Ref 37

the weldment region and compressive residual stress in the surrounding base metal region (Fig 10b).

In general, a steep residual stress gradient is developed because of the steep tendency of the thermal gradient. This may, in turn, lead to hot cracking (between columnar grains) or severe center line cracking in the weld area (Ref 48). Catastrophic failures of welded bridges and all-welded ships are mostly attributed to the existence of large and dangerous tensile residual stress in them (Ref 49).

The grinding step in manufacturing is important, since it is always utilized to produce the finished surface. It has been shown that gentle surface grinding, using a soft sharp wheel and slow downfeed, produces compressive residual stress at the surface, whereas conventional (normal practice) and abrasive grinding result in surface tensile stresses of very high magnitude (Fig l l) (Ref 22, 50). However, the

400 (58)

Distance from surface, in. 0.08 0.16 0.24 0.32 0.40 500 gm Knoop test 3 mm case I

I I Is ss

/ Hardness

2 4 6 8 10 Distance from surface, mm

8O

o r~ 60 -r-

40 ~

20 ,~

u,l

0 12

A

200 (29) #_

0

-200 (-29) tr

gentle grinding method is expensive from the viewpoint of operating time and wear of the wheel.

As a result of temperature gradient during cooling, castings develop compressive stresses at the surface and tensile stresses in the interior (Ref 22). However, transient temperature gradient and phase transformation occurring during the early stages of solidifica- tion and cooling of continuous steel castings in the mold may give rise to the development of harmful residual stresses leading to the formation of cracks (Ref 51).

Chemical processes such as electroplat- ing, scale formation, and corrosion of metals can produce residual stresses due to coherency strains arising from the matching tendency of crystal structures of the outer surface product with the crystal structure of the adjacent layer (Ref 22). Residual stresses are also introduced when heat-treated parts are subjected to successive heating and cooling cycles during service conditions.

Residual Stress in the Heat-Treated Non- ferrous Alloys. In nonferrous alloys, notably age-hardenable aluminum alloys, copper- beryllium alloys, certain nickel-base super-

(a)

Residual stress Compression Tension

i i i i i i i i i i I i i ~

Yl

(b)

Fig 10 (a) The transverse shrinkage occurring in butt weldments. (b) Longitudinal residual

stress patterns in the weldment and surrounding regions. This also shows longitudinal shrinkage in a butt weld. Source: Ref 47

-400 (-58)

F ig 9 A typical hardriess and residual stress profile in induction-hardened (to 3 ram, or 0.12 in.,

case depth) and tempered (at 260 C, or 500 F) 1045 steel. Source: Ref 44

-- 50 /xA z, /x A 0

- -50

- -100

-150 ~

--200 "~ (b

FeB - -250 rr Fe2B /x Ferrite - -300

0.10 0.15

Distance from surface, mm

Residual stress distribution of FeB and Fe2B Fig 8 layers in borided steel processed at 900 C (1650 F). Source: Ref 18, 43

alloys, and so on, a significant amount of thermal stress is generated during quenching prior to precipitation hardening. The quenching process in this condition does not invariably involve a phase change; rath- er, this is confined to the postquenching aging treatment. In other nonferrous alloys such as uranium and titanium alloys, the final structural condition is not obtained by a slow cool.

When high-strength titanium alloy is quenched from a solution annealing temperature of 850 to 1000 C (1560 to 1830 F), it develops large residual stress caused by poor thermal conductivity of titanium leading to high-temperature gradient. This problem can, however, be avoided by stress- relief annealing at 650 to 700 C (1200 to 1290 F), which produces a slight reduction in mechanical properties. When a high- strength aluminum age-hardening alloy is rapidly quenched from the solution temper-

Depth below surface, mil 3.15 6.3 9.45 12.6

800 120

600 A

g_g ~: 400

"~ 200

r,~ o

~. -200 E o

-400 0

/\.

\ \

- 9O

- 60 -- Abrasive

% 30

Conventional

0 -%

- Gentle ,-, . . . . - -30

-60 80 160 240 320

Depth below surface, pm

L~

Residual stress distribution after gentle, con- F i g 11 ventional, and abrasive grinding of hardened 4340 steel. Source: Ref 22


Table 6 A compiled summary of the maximum residual stresses in surface heat-treated steels

Residual stress (longitudinal)

Steel Heat treatment MPa ksi

832M13 (type) Carburized at 970 C (1780 F) to 1 mm (0.04 in.) case with 0.8% surface carbon

Direct-quenched 280 Direct-quenched, - 80 C ( - 110 F) subzero treatment 340 Direct-quenched, - 90 C ( -130 F) subzero treatment, 200

tempered 805A20 Carburized and quenched 240-340(a) 805A20 Carburized to 1.1-1.5 mm (0.043-0.06 in.) case at 920 C 190-230

(1690 F), direct oil quench, no temper 805A ! 7 400 805A17 Carburized to 1.1-1.5 mm (0.043-0.06 in.) case at 920 C 150-200

(1690 F), direct oil quench, tempered 150 C (300 F) 897M39 Nitrided to case depth of about 0.5 mm (0.02 in.) 400--600 905M39 800-1000 Cold-rolled steel Induction hardened, untempered 1000

Induction hardened, tempered 200 C (390 F) 650 Induction hardened, tempered 300 C (570 F) 350 Induction hardened, tempered 400 C (750 F) 170

(a) Immediately subsurface, that is. 0.05 mm (0.002 in.). Source: Ref 29

40.5 49.0 29.0

35.0--49.0 27.5-33.5

58 22-29

58.0-87.0 116.0-145.0

145.0 94.0 51 24.5

ature, high thermal and residual stresses are induced due to high coefficient of expansion of aluminum. Uphill quenching from liquid nitrogen temperature ( - 196 C, or - 320 F) in a steam blast alleviates this problem. This induces stresses opposite in sign to those developed on water quenching from the solutionizing and cancels out their effect. This is followed by aging of the alloy in the conventional manner (Ref 29).

Fast polyalkylene glycol (PAG) quenching of solution-treated aluminum alloys tends to reduce residual stress levels because of its more uniform heat extraction rate (thermal shock is smaller, and thereby machining is less likely to produce further distortion), thereby helping solve major and long-standing distortion problems among aluminum workpieces (Ref 52).

Control of Residual Stresses in Heat-Treated Parts

Table 6 lists some typical values of maximum residual stresses developed in the surface-hardened steels that have been reported in the literature (Ref 29). It is worth noting that there is a marked influence of tempering on the residual stress level. Tem- pering must be accomplished at about 150 C (300 F) to maintain 50 to 60% retention of the residual stress level obtained after quenching because a higher tempering temperature greatly reduces surface compressive stresses. However, a higher stress- relief temperature ( -600 C, or 1110 F) is used for mechanically deformed components (for example, hot-rolled bars) or components with tensile surface residual stresses. Alternatively, serious residual tensile stresses may be avoided effectively by gentle grinding of the surface.

Measurement of Residual Stresses There are two methods of measuring re-

sidual stresses: the destructive method, also

called the dissection method, and the nondestructive methods comprising mainly x-ray diffraction, neutron diffraction, ultrasonic, and magnetic methods.

Destructive (or Dissection) Method. This method is old but reasonably accurate, practically nondestructive, uses well-established methods, and can be employed in confined situations at site (Ref 53). Howev- er, it is tedious, time consuming, and expensive (Ref 54). The other drawbacks are the destructive, or at best semidestructive na- ture of the method, and its ability to measure only the macroresidual stresses. The hole-drilling method is used extensively for measuring residual stresses, which depends on the dissection approach. It consists of the mounting of strain gages or a three- element strain-gage rosette on the surface and measurement of strains. Then a rigidly guided milling cutter is used to drill a small, straight, circular, perpendicular, and fiat- bottomed hole not exceeding 3.2 mm (0.125 in.) at the center of the rosette and into the surface of the component being analyzed. Strain redistribution occurring at the surface in the surrounding area of the hole (resulting from the residual stress relief) is then measured with the previously installed strain gages. The residual stress is calculated at a large number of points in a surface from the strain measurements using the well-established method (Ref 22, 28). To minimize the introduction of spurious strains by the grinding operation, the rate of metal removal should be less than 3.125 x 10 -4 m/s (1.23 10 -2 in./s), and readings are recorded after 15 min of the end of the grinding process to ensure that any heat generated has been dissipated (Ref 55).

Nondestructive Methods. The main diffi- culty with the nondestructive methods is that measurements of crystallographic lattice parameters, ultrasonic velocities, or magnetiza- tion changes are made that are indirectly

related to the residual stress. The above quantities are usually dependent on the stress and material parameters (such as metallurgical textures), which are difficult to quantify (Ref 54, 56).

The x-ray diffraction method is the well- established technique for measuring both macro- and microresidual stress nondestruc- tively. In most instances, the x-ray diffraction method has been employed to provide quan- titative values for residual stress profiles in surface or fully hardened components (Ref 57). This technique depends on the determination of lattice strains and the stress-induced differences in the lattice spacing. Macroresid- ual strain is measured from the shift of diffraction lines in the peak position using the so-called nonlinear SinZC method from which residual stress is calculated (Ref 57). For the measurement of microstrain the Voigt single- line method is applied (Ref 58). Precision in lattice strain measurement of the order of 0.2% is possible.

Portable x-ray diffraction equipment is now commercially available in various forms that allow stress measurement to be made very quickly (ranging from 4 to 30 s). The main drawbacks are that it cannot be applied to noncrystalline materials such as plastics, and it is only capable of measuring residual stresses of materials very close to the surface under examination. That is, the measurement is purely surface related (a depth of 0.01 mm, or 0.4 mil, is commonly quoted) (Ref 59).

Neutron radiography or diffraction, used for polycrystalline materials, has a much deeper penetration than x-rays, but has major safety problems and the disadvantage of being nonportable.

Ultrasonic method for evaluating residual stress involves ultrasonic stress birefrin- gence or sonoelasticity; this depends upon the linear variation of the velocities of sound in a body (that is, ultrasonic waves) with the stress. This method has the potential for greater capability, versatility, and usefulness in the future (Ref 53, 56). How- ever, this has the disadvantage, in common with the magnetic methods, that it requires transducers shaped to match the surface being inspected (Ref 60).

The magnetic method is based on the stress dependence of the Barkhausen noise amplitude. Each time an alternating magnetic field induced in a ferromagnetic material is reversed, it generates a burst of Barkhausen noise. The peak amplitude of the burst, as determined with an inductive coil near the surface of the component material, varies with the surface stress level. Since Barkhausen noise depends on composition, texture, and work hardening, it is necessary in each application to use calibrated standard (reference) samples with the same processing history and composition as the component being analyzed. This method is used to measure residual


stresses well below the yield strength of the ferromagnetic materials. This method is rapid, and the measurements are made with the commercially available portable equipment. However, this method is limited to only ferromagnetic materials (Ref 56).

Thermal evaluation for residual stress analysis (TERSA) is a new nondestructive method that is in an experimental stage. It has the advantage that it is completely independent, remote, and noncontacting. It consists of merely directing a controlled amount of energy from a laser energy source into the volume of the material being inspected and then mak- ing a precise determination of changes in the resulting temperature rise by infrared radiom- etry. However, the working instrument will also require some form of display to enable visual examination to be made of any high- stressed regions (Ref 60).

Quench Cracking Anything that produces excessive

quenching stress is the basic cause of cracking. Quench cracking is mostly intergranular, and its formation may be related to some of the same factors that cause intergranular fracture in overheated and burned steels. The main reasons for cracking in heat treatment are: part design, steel grades, part defects, heat-treating practice, and tempering practice (Ref 61).

Part Design. Features such as sharp cor- ners, the number, location, and size of holes, deep keyways, splines, and abrupt changes in section thickness within a part (that is, badly unbalanced section) enhance the crack formation because while the one (thin) area is cooling quickly in the quenchant, the other (thick) area immediately adjacent to it is cooling very slowly. One solution to this problem is to change the material so that a less drastic quenchant (for example, oil) can be employed. An alternate solution is to prequench, that is, to cool it prior to the rest of the part. This will produce an interior of the hole or keyway that is residually stressed in compression, which is always desirable for better fatigue properties (Ref 61). The third solution is a design change, and the fourth is to use a milder quenchant.

Steel Grades. Sometimes this can be checked by means of a spark test, whereas at other times a chemical analysis must be made. In general, the carbon content of steel should not exceed the required level; other- wise, the risk of cracking will increase. The suggested average carbon contents for water, brine, and caustic quenching are given below:

Method Shape Carbon, %

Induction hardening Complex 0.33 Simple 0.50

Furnace hardening Complex 0.30 Simple 0.35 Very simple, such

as bar 0.40

A decrease in carbon content from 0.72 to 0.61% has been shown to slightly increase the thermal crack resistance of rim- quenched railroad wheels (Ref 62).

Because of segregation of carbon and alloying elements, some steels are more prone than others to quench cracking. Among these steels, 4140H, 4145H, 4150H, and 1345H appear to be the worst. A good option is to replace the 4100 series with the 8600 series. An additional disadvantage with the use of 1345H steel is the manganese floating effect, which leads to very high manganese content in the steel rolled from the last ingot in the same heat. Simi- larly, dirty steels (that is, steels with more than 0.05% S, for example, AISI 1141 and 1144) are more susceptible to cracking than the low-sulfur grades. The reasons for this are that they are more segregated in alloying elements, the surface of this hot-rolled high- sulfur steel has a greater tendency to form seams, which act as stress raisers during quenching, and they are usually coarse grained (for better machinability), which increases brittleness and therefore promotes cracking. If these high-sulfur grades are replaced by calcium-treated steels or cold-finished leaded steels, this problem can be obviated (Ref 61).

Part Defects. Surface defect or weakness in the material may also cause cracking, for example, deep surface seams or nonmetallic stringers in both hot-rolled and cold-finished bars. Other defects are inclusions, stamp marks, and so forth. For large-seam depths, it is advisable to use turned bars or even magnetic particle inspection. The forging defects in small forgings, such as seams, laps, flash line, or shearing crack, as well as in heavy forgings, such as hydrogen flakes and internal ruptures, aggravate cracking. Similarly, some casting defects, for example, in water-cooled castings, promote cracking (Ref 50).

Heat-Treating Practice. Higher austenitizing temperatures increase the tendency toward quench cracking. Similarly, steels with coarser grain size are more prone to cracks than fine-grain steels because the latter possess more grain-boundary area to stop the movements of cracks, and grain boundaries help to absorb and redistribute residual stresses. An outstanding contribu- tor to severe cracking is improper heat- treating practice, for example, nonuniform heating and nonuniform cooling of the component involved in the heat-treatment cycle. It is a good heat-treating practice to anneal alloy steels prior to the hardening treatment (or any other high-temperature treatment, for example, forging, welding, and so forth) because this produces grain- refined microstructure and relieves stresses (Ref 63).

Water-Hardening Steel. The water-hardening steels are most susceptible to cracks if they are not handled properly. Soft spots

( Typical appearance of thumbnail check as

Fig 1 2 soft spot on chipping chisel. Source: Ref 64

are most likely to occur in the water-hardening steels, especially where the tool is grabbed with tongs for quenching. Normal- ly the cleaned surface shows adequate hardening and the scaled surface insufficient hardening, which can be examined with a file. Soft spots may occur from the use of fresh water, or water contaminated with oil or soap. Most large tools emerging from hardening operations contain some soft spots. However, accidental soft spots in the wrong place should be investigated, and steps must be taken to eliminate them.

Figure 12 shows the typical appearance of a thumbnail check as soft spot on chipping chisels, which occurs on the bit near the cutting edge. The cracks enclosing the soft spots should be avoided by switching to brine quench (Ref 64).

Air-Hardening Steel. Similarly, when air hardening steels are improperly handled, they are likely to crack. For example, avoidance of tempering treatment or use of oil quenching in air-hardening steel can lead to cracking. However, the common practice in the treatment of air-hardening steels is initially to quench in oil until "black" (about 540 C, or 1000 F), followed by air cooling to 65 C (150 F) prior to tempering. As compared to air cooling right from the quenching temperature, this practice is totally safe and minimizes the formation of scale.

Polymer quenchants have found well-established use in the quenching of solution- treated aluminum alloys, hardening of plain carbon steels with less than 0.6% C, spring steels, boron steels, hardenable stainless steels, and all carburizing and alloy steels with section thickness greater than about 50 mm (2 in.), through-hardening and carburizing steel parts, and induction and flame- hardening treatments because of their nu- merous beneficial effects, including elimination of soft spots, distortion, and cracking problems associated with trace


Fig 13 Microcracking in a Ni-Cr steel. Source: Ref 67

water contamination in quenching oils (Ref 65).

Agitation is an important parameter in polymer quenching applications both to ensure a uniform polymer film around the quench part and to provide a uniform heat extraction from the hot part to the adjacent area of quenchant by preventing a buildup of heat in the quench region.

Salt bath cooling of induction-hardened complex-shaped cast iron parts reduces danger of cracking, which is usually expe- rienced when air cooling followed by hot- water quenching is used (Ref 66).

Decarburized Steel. Decarburization usually arises from insufficient protection as a result of plant failure (for example, defective furnace or container seals, defective valves), poor process control (for example, insufficient atmosphere-monitoring equipment, poor supervision), or the existence of decarburizing agents in the furnace atmosphere (for example, CO2, water vapor, and Hz in the Endogas (Ref 61, 67).

A partially decarburized surface on the part occurring during tool hardening also contributes to cracking because martensite transformation is completed therein well before the formation of martensite in the core. Decarburized surface on the tools has reduced hardness, which will lead to prema- ture wear and scuffing. Partial decarburiza-

4.7 ~m

tion must be avoided, especially on all deep- hardening steels, either by providing some type of protective atmosphere during the heating operation, stock removal by grinding, or carbon restoration process. In addition to protective atmosphere, salt baths, inert packs, or vacuum furnaces may be used to obtain the desired surface chemistry on the tools or dies. The fact that the better and more consistent performance of the tools is observed after regrinding reveals the existence of partial decarburization remaining.

Carburized Alloy Steel. Two types of peculiar cracking phenomena prevail in the carburized and hardened case of the carburized alloy steels: microcracking and tip cracking. Microcracking of quenched steels are small cracks appearing across or alongside martensite plate (Fig 13) (Ref 67) and the prior austenite grain boundaries (Ref 68). They form mostly on those quenched steel parts that contain chromium and/or molybdenum as the major alloying elements with or without nickel content and where the hardening is done by direct quenching.

Microcracks are observed mostly in coarse-grained structures, such as large martensite plates. This is presumably because of more impingements of the larger plates of martensite by other large plates.

Another cause of microcracking is the increased carbon content of martensite (that is, increased hardenability), which is a function of austenitizing temperature and/or time (Ref 67). This finding was established for 8620H steel, which has a higher austenitizing temperature prior to quenching where there is a greater tendency to micro- crack (Ref 69). This problem can be avoided by selecting a steel with less hardenability (that is, with less austenitizing temperature). Another solution is to change the heat-treating cycle to carburizing, slow cooling to black temperature, reheating to, for example, 815 or 845 C (1500 or 1550 F), and quenching (Ref 61). Microcracking in case-hardened surfaces may be aggravated by the existence of hydrogen, which tends to absorb during carburizing. However, this hydrogen-enhanced microcracking can be eliminated by tempering the carburized parts at 150 C (300 F) immediately after quenching. Tempering exhibits an additional beneficial effect in that it has the ability to heal the microcracks due to the volume changes and associated plastic flow that develop during the first stage of tempering (Ref 70). No adverse report on the influence of microcracks on the mechanical properties has been noted; however, the control- ling factors should be varied so as to keep the incidence of microcracks to a minimum (Ref 67).

Tip cracking refers to the cracking that appears in the teeth of carburized and quenched gears and runs partly or fully to the ends of the teeth in a direction parallel to the axis of the part. Many heat treaters have solved this problem to a great extent by decreasing the carbon content and case depth to the minimum acceptable design level or by copper plating the outer diameter of the gear blank prior to hobbing (Ref 66).

Nitrided Steels. The nitrided cases are very brittle. Consequently, cracking may occur in service prior to realizing any im- proved wear and galling resistance. This can be avoided by a proper tool design, for example, incorporating all section .changes with a minimum radius of 3 mm (0.125 in.).

Tempering Practice. The longer the time the steel is kept at a temperature between room temperature and 100 C (212 F) after the complete transformation of martensite in the core, the more likely the occurrence of quench cracking. This arises from the volumetric expansion caused by isothermal transformation of retained austenite into martensite.

There are two tempering practices that lead to cracking problems: tempering too soon after quenching, that is, before the steel parts have transformed to martensite in hardening, and skin tempering, usually observed in heavy sections (=>50 mm, or 2 in., thick in plates and >75 mm, or 3 in., in diameter in round bars).


It is the normal practice to temper immediately after the quenching operations. In this case, some restraint must be exercised, especially for large sections (>75 mm, or 3 in.) in deep-hardening alloy steels. The rea- son is that the core has not yet completed its transformation to martensite with the expansion, whereas the surface and/or projec- tions, such as flanges, begin to temper with shrinkage. This simultaneous volume change produces radial cracks. This problem can become severe if rapid heating practice (for example, induction, flame, lead, or molten salt bath) is used for tempering. Therefore, very large and very intri- cate tool steel parts should be removed from the quenching medium, and tempering should be started while they are slightly warm to hold comfortably in the bare hands ( -60 C, or 140 F).

Skin tempering occurs in heavy section parts when the final hardness is >360 HB. This is due to insufficient tempering time and is usually determined when the surface hardness falls by 5 or more HRC points from the core hardness. This cracking often occurs several hours after the component has cooled from the tempering temperature and often runs through the entire cross section. This problem can be removed by retempering for 3 h at the original tempering temperature, which is associated with a change in hardness of 2 HRC points maximum (Ref 61).

Distortion in Heat Treatment

Distortion can be defined as an irreversible and usually unpredictable dimensional change in the component during processing from heat treatment and from temperature variations and loading in service. The term dimensional change is used to denote changes in both size and shape (Ref71). The heat-treatment distortion is therefore a term often used by engineers to describe an un- controlled movement that has occurred in a component as a result of heat-treatment operation (Ref 72). Although it is recognized as one of the most difficult and trou- blesome problems confronting the heat treater and the heat-treatment industries on a daily basis, it is only in the simplest thermal heat-treatment methods that the mechanism of distortion is understood. Changes in size and shape of tool-steel parts may be either reversible or irreversible. Reversible changes, which are produced by applying stress in the elastic range or by temperature variation, neither induce stresses above the elastic limit nor cause changes in the metallurgical structure. In this situation, the initial dimensional values can be restored to their original state of stress or temperature.

Irreversible changes in size and shape of tool-steel parts are those that are caused by stresses in excess of the elastic limit or by

changes in the metallurgical structure (for example, phase changes). These dimensional changes sometimes can be corrected by mechanical processing to remove extra and unwanted material or to redistribute residual stresses or by heat treatment (annealing, tempering, or cold treatment).

When heat-treated parts suffer from distortion beyond the permissible limits, it may lead to scrapping of the article, rendering it useless for the service for which it was intended, or it may require necessary cor- rection. Allowable distortion limits vary to a large extent, depending on service applications; in cases where very little distortion can be tolerated, specially desired tool steels are used. These steels possess metallurgical characteristics that minimize distortion.

Types of Distortion Distortion is a general term that involves

all irreversible dimensional change produced during heat-treatment operations. This can be classified into two categories: size distortion, which is the net change in specific volume between the parent and transformation product produced by phase transformation without a change in geometrical form, and shape distortion or warpage, which is a change in geometrical form or shape and is revealed by changes of curva- ture or curving, bending, twisting, and/or nonsymmetrical dimensional change without any volume change (Ref 72, 73). Usually both types of distortion occur during a heat-treatment cycle.

Dimensional Changes Caused by Changes in Metallurgical Structure during Heat Treat- ment. Various dimensional changes produced by a change in metallurgical structure during the heat-treatment cycle of tool steels are described below (Ref 74).

Heating (Austenitizing). When annealed steel is heated from room temperature, thermal expansion occurs continuously up to Ac~, where the steel contracts as it transforms from body-centered cubic (bcc) ferrite to face-centered cubic (fcc) austenite. The extent of decrease in volumetric contraction is related to the increased carbon content in the steel composition (Table 4). Further heating expands the newly formed austenite.

Hardening. When austenite is cooled quickly, martensite forms; at intermediate cooling rates, bainite forms; and at slow cooling rates, pearlite precipitates. In all these transformation sequences, the magnitude of expansion increases with the decrease in carbon content in the austenite (Table 4). The volume increase is maximum when austenite transforms to martensite, intermediate with lower bainite, and is least with upper bainite and pearlite (Table 4). The volume increases associated with the transformation of austenite to martensite in 1 and 1.5% carbon steels are 4.1 and 3.84%,

respectively; the volume increases involved in the transformation of austenite to pearlite in the same steels are 2.4 and 1.33%, respectively. Such volume increases are less in alloy steels and least in 2C-12Cr and A10 tool steels. It should be noted that plastic deformation (or strain) occurs during such transformations at stresses that are lower than the yield stress for the phases present (Ref 75). The occurrence of this plastic deformation, called the transformation plas- ticity effect, influences the development of stresses during the hardening of steel parts (Ref 76). During quenching from the austenite range, the steel contracts until the M~ temperature is reached, then expands during martensitic transformation; finally, thermal contraction occurs on further cooling to room temperature. As the hardening temperature increases, a greater amount of carbide goes into solution; consequently, both the grain size and the amount of retained austenite are increased. This also increases the hardenability of steel.

More trouble with distortion comes from the quenching or hardening operation than during heating for hardening, in which the faster the cooling rate (that is, the more severe the quenching), the greater the danger of distortion. When the milder quenchants are used, the extent of distortion is lessened. The severity of quenching thus influences the distortion of components.

The dependence of volume increase, particularly in tools of different dimensions, on grain size (or hardenability) is another important factor. Variations in volume during quenching of a fine-grained shallow-hardening steel in all but small sections is less than a coarse-grained deep-hardening steel of the same composition.

Tempering. There is a certain correlation between the tempering temperature and volume change. Tempering reduces the volume of martensite but not adequately enough to equalize completely the prior volume increase as a result of martensitic transformation unless the components are completely softened. In low-alloy and plain (medium- and high-) carbon steels, during the first and third stages of tempering, a decrease in volume occurs that is associated with the decomposition of: high-carbon martensite into low-carbon martensite plus ~-carbide in the former stage, and aggregate of low-carbon martensite and t-carbide into ferrite plus cementite in the latter stage. In the second stage, however, an increase in volume takes place (due to the decomposition of retained austenite into bainite) that tends to compensate for the early volume reduction. As the tempering temperature is increased further toward the A~, more pro- nounced volume reduction occurs. In some highly alloyed tool-steel compositions, the volume changes during martensite formation are less striking because of the large proportion of retained austenite and the

Defects and

Table 7 Typical volume percentages of microconstituents existing in four different tool steels after their standard hardening treatments

Retained Undissolved As-quenched Martensite, austenite, carbides,

Steel Hardening treatment hardness, HRC vol% vol% vol%

W1 790 C (1450 F), 30 rain; WQ 67.0 L3 845 C (1550 F), 30 min; OQ 66.5 M2 1225 C (2235 F), 6 rain; OQ 64 D2 1040 C (1900 F), 30 rain; AC 62

Note: WQ, water quenched; OQ, oil quenched; AC, air cooled.

88.5 9 2.5 90 7 3.0 71.5 20 8.5 45 40 15

resistance to tempering of alloy-rich martensite. These hardened steels show sharp increases both in hardness and volume between 500 and 600 C (930 and 1110 F) owing to the precipitation of very finely dispersed alloy carbides from the retained austenite. This produces a depleted matrix in alloy content, raising the M~ temperature of retained austenite. During cooling down from the tempering temperature, further transformation of retained austenite into martensite will occur with an additional increase in volume.

Size Distortion. Table 7 shows the typical volume percentages of microconstituents present in four different tool steels after their standard hardening treatments. Typi- cal dimensional changes during hardening and tempering of several tool steels are given in Table 8. It is apparent here that some steels such as M3 and M41 high-speed steels show appreciable increase in size of about 0.2% after hardening and tempering between 540 and 595 C (1000 and 1100 F) to produce complete secondary hardening. Other types, such as A10, expand very little when hardened and tempered over the entire temperature range up to 595 C (1100 F). Excessive size changes in oil-hardening nonshrinkable tool steel is usually caused by lack of stress relief (when necessary), and hardening and/or tempering at the in- correct temperature. The golden rule is to learn to be suspicious of tools that are seriously off size in only one dimension. It is further noted that alloying addition in steels brings about a change in the specific volume of many microconstituents, but to a

lesser extent than carbon (Ref 77). This table provides comparative data on size distortion in a variety of steels; however, this information cannot be used alone to predict shape distortion factor.

Shape Distortion or Warpage. This is sometimes called straightness or angularity change. It is found particularly in nonsymmetrical components during heat treatment. From the practical viewpoints, warpage in water- or oil-hardening steels is normally of greater magnitude than is size distortion and is more of a problem because it is usually not predictable. This is caused by the sum effect of more than one of these factors:

Rapid heating (or overheating), drastic (or careless) quenching, or nonuniform heating and cooling causes severe shape distortion. Slow heating as well as pre- heating of the parts prior to heating to the austenitizing temperature yields the most satisfactory result. Rapid quenching produces thermal and mechanical stresses associated with the martensitic transformation. In the case of low- and high-hardenability steels, respectively, this problem becomes severe or very small

Residual stresses present in the component before heat treating. These arise from machining, grinding, straightening, welding, casting, spinning, forging, and rolling operations, which will also furnish a marked contribution to the shape change (Ref 78)

Applied stress causing plastic deformation. Sagging and creep of the compo-

Distortion in Heat-Treated Parts / 613

nents occur during heat treatment as a result of improper support of components or warped hearth in the hardening furnace. Hence, large, long, and complex- shaped parts must be properly supported at critical positions to avoid sagging or preferably are hung with the long axis on the vertical

Nonuniform agitation/quenching or nonuniform circulation of quenchant around a part results in an assortment of cooling rates that creates shape distortion (Ref 79). Uneven hardening, with the formation of soft spots, increases warpage. Similarly, an increase in case depth, particularly uneven case depths in case-hardening steels, increases warpage on quenching (Ref 80)

Tight (that is, thin and highly adherent) scale and decarburization, at least in certain areas. Tight scale is usually a problem encountered in forgings hardened from direct-fired gas furnaces having high-pressure burners. Quenching in areas with tight scale is extremely retarded compared to the areas where the scale comes off. This produces soft spots, and, in some cases, severe unpredicted distortion. Some heat treaters coat the components with a scale-loosening chemical pri- or to their entry into the furnace (Ref 79). Similarly, the areas beneath the decarburized surface do not harden as completely as the areas below the nondecarburized surface. The decarburized layer also varies in depth and produces an inconsistent softer region as compared to the region with full carbon. All these factors can cause a condition of unbalanced stresses with resultant distortion (Ref 79)

Long parts with small cross sections (>L = 5d for water quenching, > L = 8d for oil quenching, and > L = 10d for austempering, where L is the length of the part, and d is its diameter or thickness)

Thin parts with larger areas (>A = 50t, where A is the area of the part, and t is its thickness)

Unevenness of, or greater variation in, section

Table 8 Typical dimensional changes during hardening and tempering of several tool steels Hardening treatment

Tool Temperature Quenching steel ~U- 1 medium

Total change in linear dimensions

after quenching, %

Total change in linear dimensions~ %, after tempering at

150 *C 205 *C 260 *C 315 *C 370 *C 300 *F 400 *F 500 *F 600 *F 700 *F

425 *C 480 *C 510 *C 540 *C 565 *C 595 *C 800 *F 900 *F 950 oF 1000 oF 1050 *F 1100 *F

Ol 815 1500 Oil 0.22 OI 790 1450 Oil 0.18 06 790 1450 Oil 0.12 A2 955 1750 Air 0.09 A10 790 1450 Air 0.04 D2 11)10 1850 Air 0.06 D3 955 1750 Oil 0.07 D4 1040 1900 Air 0.07 D5 1010 1850 Air 0.07 H I I 1010 1850 Air 0.11 HI3 1010 1850 Air -0.01 M2 1210 2210 Oil -0 .02 M41 1210 2210 Oil -0 .16

0.17 0.16 0.18 0.09 0.12 0.13 - 0.07 0.10 0.14 0.10 0.06 0.06 0.08 0.07 0.00 0.00 0.08 0.08 0.03 0.03 0.02 0.00 0.04 0.02 0.01 -0 .02 0.03 0.01 -0.01 -0 .03 0.03 0.02 0.01 0.00 0.06 0.07 0.08 0.08

0.00 -0 .05 -0 .06 0.05 0.04 0.01 0.01 0.02 . . . . 0.01 -0 .02

- 0 . 4 -0.03 ' 0.3 0.03 0.3 0.01 . . . . . 0.00

-0 .06 -0 .17

-0 .07 0.06 0.01 0.06

0.05 0.05 0.12 0.06 0.10 0.08

0.14 0,21

0.02

0.16 0.23


Examples of Distortion

Ring Die. Quenching of ring die through the bore produces the reduction in bore diameter as a result of formation of martens- ire, assoc ia ted with the increased volume. In other words, metal in the bore is upset by shrinkage of the surrounding metal and is short when it cools (Ref 24). However , allover quenching causes the outside diameter to increase and the bore diameter to increase or decrease, depending upon precise dimensions of the part. When the outside diameter of the steel part is induction- or f lame-hardened (with water quench), it causes the part to shrink in outer diameter (Ref 63). These are the examples of the effect of mode of quenching on distortion (Ref 81).

Thin die (with respect to wall thickness) is likely to increase in bore diameter, decrease in outside diameter, and decrease in thickness when the faces are hardened. If the die has a very small hole, insufficient quenching of the bore may enlarge the hole diameter because the body of die moves with the outside hardened portion.

Bore of Finished Gear. Similarly, the bore of a finished gear might turn oval or change to such an extent that the shaft cannot be fitted by the allowances that have been provided. Even a simple shape such as a diaphragm or orifice plate may, after heat treatment, lose its flatness in such a way that it may become unusable.

Production of Long Pins. In the case of the production of long pins (250 mm long x 6 mm diameter, or 10 V4 in.) made from medium-alloy steel, it was found, after conventional hardening, that when mounted between centers, the maximum swing was over 5 mm (0.20 in.). However , the camber could be reduced to within acceptable limits by martempering, intense or press quenching.

Hardening and Annealing of Long Bar. When a 1% carbon steel bar, 300 mm long (or more) 25 mm diameter (12 in. long, or more, 1 in. diameter), is water quenched vertically from 780 C (1435 F), the bar increases both in diameter and volume but decreases in length. When such bars are annealed or austenitized, they will sag badly between the widely spaced supports. Hence, they should be supported along their entire length in order to avoid distortion.

Hardening of Half-Round Files. Files are usually made from hypereutectoid steel containing 0.5% chromium. Files are heated to 760 C (1400 F) in an electric furnace after being surface coated with powdered wheat, charcoal , and ferrocyanide to pre- vent decarburization. They are then quenched vertically in a water tank. On their removal from the tank, the files appear like the proverbial dog 's tail. The flat side has curved down, the camber becomes ex-

cessive, and the files can no longer be used in service. One practical solution is to give the files a reverse camber prior to quenching. The dead fiat files could, however, be made possible, and the judgment with re- gard to the actual camber needed depends upon the length and the slenderness of the recur files (Ref 82).

Similarly, when a long slender shear knife is heat treated, it tends to curve like a dog 's tail, unless special precautions are taken.

Hardening of Chisels (Ref 63). Chisels about 460 mm (18 in.) long and made from 13 mm (0.5 in.) AISI 6150 bar steel are austenitized at 900 C (1650 F) for 1.5 h and quenched in oil at 180 C (360 F) by standing in the vertical position with chisel point down in special baskets that allow stacking of two 13 mm (0.5 in.) round chisels per 650 mm 2 (1 infl) hole. Subsequently, hardened chisels are tempered between 205 and 215 C (400 and 420 F) for 1.5 h. These heat- treated parts show 55 to 57 HRC hardness but are warped. The reasons for this distortion are:

The portion of the bar that touches the basket cools slowly, producing uneven contraction and thermal stress

The martensite formation is delayed on the inner or abutting side of the bar, causing unequal expansion during transformation. This distortion can be eliminated or minimized by loading the parts in the screen-basket in such a way that stacking arrangement permits sufficient space between each part and by slightly decreasing the austenitizing temperature (Ref 62). Distortion can also be minimized by austempering the part, provided that the carbon content is on the high side of specification to produce the lower bainitic structure of 55 to 57 HRC. If higher yield stress is not warranted, only chisel ends need hardening and subsequent tempering (Ref 63)

Hardening of a Two-Pounder Shot. The hardness of a two-pounder shot was specified at 60 HRC on the nose and 35 HRC at the base. A differential hardening technique was performed on the shot made of a Ni-Cr-Mo steel. This technique consisted of quenching the shot in the ice-cold water by its immersion in a tank up to the shoulder, followed by drawing out the water from the tank at a stipulated rate until the water line reached the base of the nose. The final step involved withdrawing the shot from the tank when completely cold. The back end was then softened by heating in a lead bath after initial tempering. The first few shots hardened in this way were observed to split vertically across the nose. The failure was, however, avoided by withdrawal of the shot before attaining ice-cold temperature and its subsequent immersion in warm water (Ref 82).

Hardening of a Burnishing Wheel. In the manufacture of railway axles, the gearing

surface on which the axle rests in the hous- ing has to be given a high burnishing polish employing a circular pressure tool that is made of !.2C-1.5Cr steel. Fo r satisfactory results, the hardness of the tool surface should be about 60 HRC. It has been found that the tool usually cracks before its withdrawal from the cold-water quenching bath. T

defects and distortion in heat-treated

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