Volume 01 - Properties and Selection Irons, Steels, and High-Performance Alloys Episode 9 pptx

5 Influence of sulfur on the transition temperature of purified iron containing 2000 ppm O, and the influence of carbon content on sulfur embrittlement.. Metalloids such as phosphorus,

Fig 30 Fatigue data (a) showing sequence effects for notched-specimen and smooth-specimen simulations

(2024-T4 aluminum, Kf = 2.0) Load histories A and B have a similar cyclic load pattern (dS2) but have slightly

different initial transients (∆S1) with either (b) a tensile leading edge (first stress peak at +∆S1 /2) or (c) a

compressive leading edge (first stress peak at -∆S1/2) The sequence effect on fatigue life (a) becomes more

pronounced as ∆S2 smaller Source: Ref 21

Ultimately, the fatigue analyst will be required to include and correlate a number of material, shape, processing, and load factors in order to identify the critical locations within a part and to describe the local stress-strain response at those critical locations The ability to anticipate pertinent factors will greatly affect the final accuracy of the life prediction

Load data gathering is one remaining topic that must be included in any discussion of fatigue Reference 21 discusses three load histories, suspension, transmission, and bracket vibration, that typify loads found in the ground vehicle industry Additionally, there are vastly different histories unique to other industries, like the so-called ground-air-ground cycle in aeronautics Without the ability to completely and accurately characterize anticipated and, occasionally, unanticipated customer use and resultant loads, the analyst will not be able to predict accurately the suitability of a new or revised design

The last several years have seen a major change in the ability to gather customer or simulated customer load data Testing methods have progressed from bulky, multichannel analogue tape recorders (where it took days or weeks before results were available) through portable frequency-modulated telemetry packages (where analysis could be performed immediately at a remote site) to hand-held packages capable of data acquisition and analysis on board the test vehicle in real time Microelectronics is further reducing size and improving reliability to the point that data can be gathered from within small, complex, moving, hostile assemblies, such as engines

References cited in this section

6 R.C Juvinall, Engineering Considerations of Stress, Strain and Strength, McGraw-Hill, 1967

7 N.E Dowling, W.R Brose, and W.K Wilson, Notched Member Fatigue Life Predictions by the Local

Strain Approach, in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed.,

Society of Automotive Engineers, 1977

8 J.A Graham, Ed., Fatigue Design Handbook, Society of Automotive Engineers, 1968

9 H.O Fuchs and R.I Stephens, Metal Fatigue in Engineering, John Wiley & Sons, 1980

10 Special Publication P-109, in Proceedings of the SAE Fatigue Conference, Society of Automotive

Engineers, 1982

11 R.C Rice, Ed., Fatigue Design Handbook, 2nd ed., Society of Automotive Engineers, 1988

12 J.T Ransom, Trans ASM, Vol 46, 1954, p 1254-1269

13 "Fatigue Properties," Technical Report, SAE J1099, Society of Automotive Engineers, 1977

14 P.H Wirsching and J.E Kempert, A Fresh Look at Fatigue, Mach Des., Vol 48 (No 12), 1976, p 120-123

15 L.E Tucker, S.D Downing, and L Camillo, Accuracy of Simplified Fatigue Prediction Methods, in

Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of Automotive

Engineers, 1977

16 R.E Peterson, Stress Concentration Design Factors, John Wiley & Sons, 1974

17 R.J Roark, Formulas for Stress and Strain, 4th ed., McGraw-Hill, 1965

18 J.M Potter, Spectrum Fatigue Life Predictions for Typical Automotive Load Histories and Materials Using

the Sequence Accountable Fatigue Analysis, in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of Automotive Engineers, 1977

19 D.V Nelson and H.O Fuchs, Predictions of Cumulative Fatigue Damage Using Condensed Load Histories,

in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of Automotive

Engineers, 1977

20 S.D Downing and D.F Socie, Simple Rainflow Counting Algorithms, Int J Fatigue, Jan 1981

21 D.F Socie, "Fatigue Life Estimation Techniques," Technical Report 145, Electro General Corporation

Fatigue Resistance of Steels

Bruce Boardman, Deere and Company, Technical Center

References

1 R.W Hertzberg, Deformation and Fracture Mechanics of Engineering Materials, John Wiley & Sons,

1976

2 D.J Wulpi, Understanding How Components Fail, American Society for Metals, 1985

3 Fatigue and Microstructure, in Proceeding of the ASM Materials Science Seminar, American Society for

6 R.C Juvinall, Engineering Considerations of Stress, Strain and Strength, McGraw-Hill, 1967

7 N.E Dowling, W.R Brose, and W.K Wilson, Notched Member Fatigue Life Predictions by the Local

Strain Approach, in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed.,

Society of Automotive Engineers, 1977

8 J.A Graham, Ed., Fatigue Design Handbook, Society of Automotive Engineers, 1968

9 H.O Fuchs and R.I Stephens, Metal Fatigue in Engineering, John Wiley & Sons, 1980

10 Special Publication P-109, in Proceedings of the SAE Fatigue Conference, Society of Automotive

Engineers, 1982

11 R.C Rice, Ed., Fatigue Design Handbook, 2nd ed., Society of Automotive Engineers, 1988

12 J.T Ransom, Trans ASM, Vol 46, 1954, p 1254-1269

13 "Fatigue Properties," Technical Report, SAE J1099, Society of Automotive Engineers, 1977

14 P.H Wirsching and J.E Kempert, A Fresh Look at Fatigue, Mach Des., Vol 48 (No 12), 1976, p 120-123

15 L.E Tucker, S.D Downing, and L Camillo, Accuracy of Simplified Fatigue Prediction Methods, in

Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of Automotive

Engineers, 1977

16 R.E Peterson, Stress Concentration Design Factors, John Wiley & Sons, 1974

17 R.J Roark, Formulas for Stress and Strain, 4th ed., McGraw-Hill, 1965

18 J.M Potter, Spectrum Fatigue Life Predictions for Typical Automotive Load Histories and Materials

Using the Sequence Accountable Fatigue Analysis, in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of Automotive Engineers, 1977

19 D.V Nelson and H.O Fuchs, Predictions of Cumulative Fatigue Damage Using Condensed Load

Histories, in Fatigue Under Complex Loading: Analyses and Experiments, R.M Wetzel, Ed., Society of

Automotive Engineers, 1977

20 S.D Downing and D.F Socie, Simple Rainflow Counting Algorithms, Int J Fatigue, Jan 1981

21 D.F Socie, "Fatigue Life Estimation Techniques," Technical Report 145, Electro General Corporation

If the embrittlement occurs during processing at the mill, it may be detected during routine testing depending on the degree of embrittlement and the nature of the testing program The steelmaker is generally aware of the potential problems that particular grades are susceptible to and will check the various well-known parameters that can promote such problems This starts with an examination of chemical composition, such as gas contents or impurity elements that are known to cause problems For example, in the melting of ingots for pressure vessels or rotors where temper embrittlement is a potential problem, the selection of scrap for electric furnace melting is based on the scrap impurity level, and every effort is made to keep residual levels of phosphorus, antimony, tin, and arsenic as low as possible Subsequent impact testing of coupons from the forgings is used to verify the initial toughness Such information is compared to specification requirements and historical data base information to maintain quality Special melting practices, such as vacuum carbon deoxidation, have also been adopted for critical composition control Furthermore, additional knowledge is obtained by surveillance testing and by postmortems on retired forgings

However, not all potential problems can be detected, or prevented, at the mill Some arise from handling or fabrication problems For example, low-carbon sheet steels that are not aluminum killed are roller leveled at the mill prior to shipment to suppress the yield point and prevent strain-age embrittlement However, if this steel is not formed within a certain time period, the yield point will return, and discontinuous yielding may occur when the sheet is cold formed, resulting in cosmetically damaging Lüders bands on the product For this and other fabrication problems, postfabrication inspection and testing programs must be properly planned and executed

Service and environmental conditions can also cause a number of embrittlement problems Engineers working with steel components that are susceptible to such operating-induced problems must be aware of the potential problems and must establish regular inspection programs to detect problems before they become critical An excellent example of such

programs is the on-site, in situ examination of steam piping in electric power generation plants; these examinations make

extensive use of field metallography and replication to assess creep damage and predict remnant life

Embrittlement of Steels

George F Vander Voort, Carpenter Technology Corporation

Embrittlement of Iron

Before discussing the embrittlement of steels, this article will first examine the embrittlement of iron because the number

of such studies are few compared to those for steels Grain-boundary segregation of elements such as oxygen, sulfur, phosphorus, selenium, and tellurium is known to produce intergranular brittle fractures in iron at low temperatures Studies of the effects of such impurities in pure iron have been greatly aided by the development of Auger electron spectroscopy (AES)

Intercrystalline fractures in iron at low temperatures occur when the carbon content is low It has been assumed that the absence of carbon is more important than the presence of embrittling grain-boundary impurities However, impurities must be present, and the role of carbon appears to be one of a competitor for grain-boundary sites when such impurities are present

Oxygen. A study of the toughness of iron-oxygen alloys found that intergranular embrittlement and a rise in impact transition temperature in iron were produced by oxygen levels of 30 ppm or more (Ref 1) Figure 1 shows the Charpy V-notch impact transition curves from this study for a series of iron-oxygen alloys with increasing oxygen contents (from 10

to 2700 ppm) However, the sulfur levels of the heats were rather high, from 30 to 54 ppm for the alloys shown in Fig 1, except for the 0.011% O heat, which had 16 ppm S Because this work was conducted long before the introduction of AES, the true nature of the embrittling grain-boundary element(s) is open to question

Fig 1 Charpy V-notch impact energy curves for iron-oxygen alloys with varying oxygen content Source: Ref 1

Work on high-purity iron found intergranular brittleness at low test temperatures when specimens were quenched from temperatures where carbon was in solution (Ref 2, 3) It was believed that this brittleness was due to oxygen Lowering the quench temperature reduced the embrittlement Titanium additions were not found to be helpful in reducing the intergranular embrittlement

Reference 4 reviews the study from Ref 3 and reports on similar work performed in France that found that sulfur, not oxygen, was the cause of the embrittlement Tests were done using electrolytic iron containing 35 ppm S and a purified iron containing less than 1 ppm S; each contained 2000 ppm O and varying carbon contents The purified iron was found

to be much less brittle than the electrolytic iron (Fig 2) At 10 ppm C, the purified iron was free of intergranular fracture; the electrolytic iron fractured intergranularly below the ductile-to-brittle transition temperature (DBTT)

Fig 2 Transition temperature, Tt, versus carbon content for two different high-purity irons, each containing

2000 ppm O Source: Ref 4

Reference 5, on the other hand, shows Auger analysis of intergranular fractures of the pure iron specimen from Ref 2 that contained 60 ppm S Sulfur was not detected on the intergranular fracture, while carbon, nitrogen, and oxygen were The authors did state that the oxygen peak could possibly be due to oxygen contamination after fracture in the high vacuum used for Auger work (fracture made inside the evacuated chamber)

It has been demonstrated that large variations in oxygen content have no influence on the brittleness of iron (Ref 6) The Charpy U-notch transition temperatures of electrolytic iron and high-purity iron with varying oxygen and carbon contents were determined, as in Ref 4 Figure 3 shows that the DBTTs for these two irons were constant over a wide range of oxygen contents (1 to 2000 ppm) The DBTT of the electrolytic iron (210 °C, or 410 °F) was consistently much higher than that of the high-purity iron (20 °C, or 70 °F) Also, fractures for the electrolytic iron specimens in the brittle range were fully intergranular, whereas those for the high-purity iron were by cleavage Carbon and sulfur contents were 30 and

35 ppm, respectively, for the electrolytic iron and 10 and less than 3 ppm, respectively, for the high-purity iron Manganese contents were 12 ppm for the electrolytic iron and less than 0.5 ppm for the high-purity iron, too little to tie up the sulfur completely in the electrolytic iron Figure 4 shows the results when carbon and oxygen were both varied When the carbon level was increased to 30 ppm, there was a large improvement in the DBTT, irrespective of the oxygen level This study also demonstrated that the addition of elements that form sulfide precipitates improved the DBTT

Fig 3 Ductile-to-brittle transition temperatures (from tests using Charpy U-notch specimens) as a function of

oxygen content for a decarburized electrolytic iron and a high-purity iron with 10 ppm C Source: Ref 6

Fig 4 Ductile-to-brittle transition temperatures of high-purity iron as a function of carbon content and oxygen

content Source: Ref 6

On the other hand, a study using irons with less than 2 ppm S and 0.5 ppm C and Auger analysis demonstrated oxygen grain-boundary enrichment of intergranular fractures broken by impact within the AES chamber (Ref 7) Segregation of elements such as sulfur, phosphorus, and nitrogen was not observed on the fracture, but oxygen was detected For a specimen with a bulk oxygen content of 430 at ppm, there were two monolayers of oxygen at the intergranular fracture surface Carbon was not present at the fracture surface, consistent with the very low bulk carbon content Another specimen with 235 at ppm O exhibited about 15% cleavage fracture, along with areas that were predominantly intergranular The study found that for the material with higher oxygen content, quenching specimens from temperatures

of 1103 and 1123 °C (2017 and 2053 °F), near the solubility limit, gave predominantly intergranular low-temperature fractures When the temperature was lowered from 973 to 773 °C (1783 to 1423 °F), below the solubility limit, the amount of intergranular fracture decreased Also, if the specimen was slow cooled from temperatures near the solubility limit, low-temperature fractures were fully by cleavage Therefore, the study concluded that intergranular brittle fracture

of iron was due to the oxygen in solution

Sulfur. The embrittlement of iron by sulfur has been demonstrated by several authors Prior to Auger analysis, such results were assumed, but without direct proof Auger work has, indeed, confirmed that sulfur is a potent embrittler of iron, even at bulk levels as low as 10 ppm Furthermore, the displacement of sulfur on the grain boundaries when carbon

is added has been proved by Auger analysis Low-temperature impact tests performed on three heats of relatively pure iron obtained intergranular fractures, depending on the heat treatment used (Ref 8) Sulfur contents ranged from 14 to

100 ppm, carbon contents from less than 10 to 70 ppm, and oxygen contents from 8 to 420 ppm Auger analysis revealed heavy segregation of sulfur to the grain boundaries No clear evidence of oxygen on the intergranular fractures was obtained

Other studies have shown the detrimental influence of small sulfur additions (Ref 9, 10) Figure 5, from these studies, shows the DBTT as a function of bulk sulfur content (≤60 ppm) and the beneficial influence of carbon additions (≤10 to

30 ppm) for iron containing 2000 ppm O For less than or equal to 10 ppm C, the transition temperature increased from 0

to over 600 °C (30 to >1110 °F) with increasing sulfur content However, with 25 of 30 ppm C, the influence of sulfur was small Also demonstrated was the scavenging influence of a 0.5% Al addition, which suppressed embrittlement (constant DBTT) for the full range of sulfur (≤60 ppm) tested

Fig 5 Influence of sulfur on the transition temperature of purified iron containing 2000 ppm O, and the

influence of carbon content on sulfur embrittlement Increasing carbon content has the beneficial effect of decreasing sulfur embrittlement Source: Ref 9, 10

Selenium and tellurium are similar to sulfur and are potential embrittlers of pure iron One study (Ref 10) showed that they do cause embrittlement, but to a lesser degree than sulfur (Fig 6) In this work, the carbon content was less than

10 ppm, and 2000 ppm O was present Although selenium and tellurium had less influence on the DBTT, these two elements reduced the absorbed energy values much more than did sulfur Therefore, in terms of impact energy, the elements can be placed in order of increasing influence as follows: sulfur, selenium, and tellurium Others have also reported the embrittlement of Fe-0.04% C by 0.02% Te (Ref 11)

Fig 6 Influence of sulfur, tellurium, and selenium on the transition temperature of purified iron containing up

to 10 ppm carbon and approximately 2000 ppm O Source: Ref 10

Other Impurity Elements. Phosphorus has also been reported to embrittle pure iron (Ref 12, 13) However, both of these studies used materials with a significant sulfur content, and they were performed prior to the development of Auger analysis However, radioactive tracer analysis demonstrated the segregation of phosphorus at the grain boundaries of an Fe-0.09% P alloy (Ref 12) Phosphorus was reported to be 50 times as prevalent at the grain boundaries as in the grain interiors Phosphorus does substantially increase the strength of ferrite Again, the addition of carbon was shown to reduce the influence of phosphorus on embrittlement

Researchers have also studied phosphorus segregation in pure iron (Ref 14) Again, the specimens contained a significant amount of sulfur, but mechanical properties were not determined However, a small amount of manganese was present that should precipitate the sulfur as a sulfide The carbon content was reduced to below 10 ppm Specimens were austenitized, water quenched, tempered at 850 °C (1560 °F) for 1 h, and then furnace cooled to the aging temperature Specimens were fractured within the Auger chamber Phosphorus was observed on the surface of intergranular fractures, but not on cleavage fractures Auger analysis showed that the amount of phosphorus on the intergranular fractures increased with bulk phosphorus content Also, as the aging temperature decreased, the grain-boundary phosphorus content increased, and the fracture became more intergranular When carbon was added (≤80 ppm), the grain-boundary phosphorus concentration decreased A deep-drawing steel containing 7 ppm C, 310 ppm P, and 0.36% Mn fractured intergranularly in the drawing direction, and phosphorus was detected on the grain boundaries (Ref 14) Similar steels with 14 ppm C, 80 ppm P, and 0.38% Mn did not fracture during deep drawing

A study of the embrittlement of iron by phosphorus, phosphorus and sulfur, and antimony and sulfur demonstrated that the embrittlement was different from that of temper embrittlement in that it was not reversible (Ref 15) Segregation

occurred at all temperatures in ferrite but was negligible or limited in austenite Quenching from the austenite region produced specimens that fractured by cleavage When quenched from the two-phase region, fractures did exhibit phosphorus at the grain boundaries When an Fe-0.2P alloy was furnace cooled from the austenite region, the fracture surface exhibited a layer of nearly pure phosphorus at the grain boundary with a thickness of 1 to 1.5 nm (10 to 15 Ao ) The ternary alloys containing sulfur exhibited DBTTs of about 350 °C (660 °F) The study concluded that sulfur, even at much lower concentrations than phosphorus, is a more potent embrittler of iron

Metalloids such as phosphorus, arsenic, antimony, and tin do not produce embrittlement of pure iron containing minor amounts of carbon in the same manner as sulfur, although they do in alloy steels (Ref 16) It has been demonstrated that such elements produce embrittlement of carbide/ferrite and surrounding ferrite/ferrite interfaces (Ref 17) This appears to

be a nonequilibrium segregation problem, however

The influence of tin on high-purity iron and low-carbon steel has been examined (Ref 18) Detailed chemical analyses of the pure iron specimens used were not given in the study, although it was stated that the base metal had a carbon content

of 20 ppm and an oxygen content of 400 ppm The addition of 0.5% Sn to the pure iron reduced the impact strength in the ductile region to such a degree that the absorbed energy was constant up to 70 °C (160 °F) Specimens water quenched from 650 °C (1200 °F) exhibited impact results similar to those of tin-free pure iron, while slowly cooled specimens were embrittled The addition of 0.15% C to the Fe-0.5Sn alloy did reduce the embrittling influence of tin, and the alloy had better toughness than the pure iron specimen when water quenched from 650 °C (1200 °F) The addition of 0.15% P to the Fe-0.5Sn-0.15C alloy raised the DBTT about 20 °C (36 °F) and lowered the upper-shelf energy when water quenched from 650 °C (1200 °F) The examination of fractured specimens showed a changed from transgranular cleavage to intergranular fracture as the tin content increased, particularly for the slowly cooled specimens

This survey of the influence of impurities on the embrittlement of pure iron has demonstrated that the design of experiments and the interpretation of results are difficult Many of the early studies did not recognize the significance of relatively minor amounts of sulfur in the high-purity irons used It is clear that Auger analysis is required to determine the embrittling species When sulfur is present in the absence of sulfide-forming elements, it has a dominating influence on properties and behavior The addition of carbon above about 10 ppm will reduce or eliminate embrittlement effects

References cited in this section

1 W.P Rees and B.E Hopkins, Intergranular Brittleness in Iron-Oxygen Alloys, J Iron Steel Inst., Vol 172,

Dec 1952, p 403-409

2 C.J McMahon, Jr., Intergranular Brittleness in Iron, Acta Metall., Vol 14, July 1966, p 839-845

3 J.R Rellick and C.J McMahon, Jr., The Elimination of Oxygen-Induced Intergranular Brittleness in Iron by

Addition of Scavengers, Metall Trans., Vol 1, April 1970, p 929-937

4 P Jolly, Discussion of "The Elimination of Oxygen-Induced Intergranular Brittleness in Iron by Addition of

Scavengers," Metall Trans., Vol 2, Jan 1971, p 341-342

5 J.R Rellick et al., Further Information on Oxygen Induced Intergranular Brittleness in Iron, Metall Trans.,

12 M.C Inman and H.R Tipler, Grain-Boundary Segregation of Phosphorus in an Iron-Phosphorus Alloy and

the Effect Upon Mechanical Properties, Acta Metall., Vol 6, Feb 1958, p 73-84

13 B.E Hopkins and H.R Tipler, The Effect of Phosphorus on the Tensile and Notch-Impact Properties of

High-Purity Iron and Iron-Carbon Alloys, J Iron Steel Inst., Vol 188, March 1958, p 218-237

14 H Erhart and H.J Grabke, Equilibrium Segregation of Phosphorus at Grain Boundaries of Fe-P, Fe-C-P,

Fe-Cr-P, and Fe-Cr-C-P Alloys, Met Sci., Vol 15, Sept 1981, p 401-408

15 P.V Ramasubramanian and D.F Stein, An Investigation of Grain-Boundary Embrittlement in Fe-P, Fe-P-S,

and Fe-Sb-S Alloys, Metall Trans., Vol 4, July 1973, p 1735-1742

16 C.J McMahon, Jr., Strength of Grain Boundaries in Iron-Base Alloys, in Grain Boundaries in Engineering Materials, Claitor, 1975, p 525-552

17 J.R Rellick and C.J McMahon, Jr., Intergranular Embrittlement of Iron-Carbon Alloys by Impurities,

Metall Trans., Vol 5, Nov 1974, p 2439-2450

18 C.J Thwaites and S.K Chatterjee, Effect of Tin on the Impact Behavior of Alloys Based on High-Purity

Iron and Mild Steel, J Iron Steel Inst., Vol 210, Aug 1972, p 581-587

Embrittlement of Steels

George F Vander Voort, Carpenter Technology Corporation

Embrittlement in Carbon Steels and Alloy Steels

Several forms of embrittlement can occur during thermal treatment or elevated-temperature service of carbon and alloy steels These forms of embrittlement (and the types of steel that some forms specifically affect) are:

• Blue brittleness (carbon steels)

• Quench-age embrittlement (low-carbon steels)

• Strain-age embrittlement (low-carbon steels)

• Aluminum nitride embrittlement (carbon and alloy steels

• Graphitization (carbon and alloy steels)

Blue Brittleness

Carbon steels generally exhibit an increase in strength and a reduction of ductility and toughness at temperatures around

300 °C (570 °F) Because such temperatures produce a bluish temper color on the surface of the specimen, this problem has been called blue brittleness (Ref 19, 20, 21, 22) It is generally believed that blue brittleness is an accelerated form of strain-age embrittlement Deformation in the blue-heat range followed by testing at room temperature produces an increase in strength that is greater than when the deformation is performed at ambient temperature Blue brittleness can be eliminated if elements that tie up nitrogen are added to the steel, for example, aluminum or titanium

Quench-Age Embrittlement

If a carbon steel is heated to a temperature slightly below its lower critical temperature and then quenched, it will become harder and stronger but less ductile (Ref 23, 24, 25, 26, 27, 28) This problem has been called quench aging or quench-age embrittlement The degree of embrittlement is a function of time at the aging temperature Aging at room temperature requires several weeks to reach maximum embrittlement Lowering the quenching temperature reduces the degree of embrittlement Quenching from temperatures below about 560 °C (1040 °F) does not produce quench-age embrittlement Carbon steels with a carbon content of 0.04 to 0.12% appear to be most susceptible to this problem

Quench aging is caused by the precipitation of carbide and/or nitride from solid solution One study has reviewed the quench aging of iron-nitrogen, iron-carbon, and iron-carbon-nitrogen alloys (Ref 25) Figure 7 shows the change in hardness, interparticle spacing, and practice size of precipitates in an Fe-0.02N alloy quenched from 500 °C (930 °F) and

aged at 60 °C (140 °F) up to 500 h The precipitates grew from about 30 nm (300 Ao ) in diameter after 2 h to about 80 nm (800 Ao ) in diameter after 500 h With aging, the hardness increased rapidly to about 150 HV after 10-h aging at 60 °C (140 °F) and then decreased to about 120 HV with aging to 500 h

Fig 7 Influence of aging time at 60 °C (140 °F) after quenching from 500 °C (930 °F) on the hardness, particle

size, and interparticle spacing for an Fe-0.02N alloy Source: Ref 25

Low-Carbon Steels. For low-carbon steels, quench aging is due mainly to carbon because the nitrogen level is usually too low to have a substantial effect Results for a rimming steel containing 0.03% C (Fig 8) show the increase in hardness with aging time at 60 °C (140 °F) for specimens quenched from 725 °C (1335 °F) The tensile strength increased rapidly

to a maximum value after 10 h at 60 °C (140 °F) and then decreased slowly with further aging The yield strength also increased rapidly and reached a maximum in about the same time but remained constant with continued aging The elongation and reduction of area decreased as the strength increased; they reached a minimum after 10 h and then increased somewhat with continued aging

Fig 8 Influence of aging time at 60 °C (140 °F) after quenching from 725 °C (1335 °F) on the tensile

properties of an Fe-0.03C rimming alloy (a) Tensile and yield strength (b) Elongation and reduction in area Source: Ref 25

Iron-Nitrogen and Iron-Carbon Alloys. For iron-nitrogen alloys, two types of nitrides can precipitate during quench aging:

• Face-centered cubic (fcc) Fe4N platelets form at high temperatures and generally are found at grain boundaries

• Body-centered cubic (bcc) Fe16N2 nitrides with a circular disk shape precipitate at low temperature on dislocations

This latter type of nitride causes strengthening during quench in iron-nitrogen alloys (Ref 25) Iron-carbon alloys also have two types of carbides that can form during quench aging Cementite forms at high temperatures, and a low-temperature carbide that is identical in morphology and habit to Fe16N2 and may be isomorphous with it can also form With aging, the low-temperature carbide will gradually be replaced by Fe3C (Ref 25) The phase changes during aging of iron-nitrogen and iron-carbide quench aged steels are discussed in the literature (Ref 28)

Strain-Age Embrittlement

Strain aging occurs in low-carbon steels deformed certain amounts and then aged, which produces an increase in strength and hardness but a loss in ductility (Ref 25, 26, 27, 29, 30, 31, 32, 33, 34) The degree of deformation, or cold work, is

important Generally, about a 15% reduction in thickness provides the maximum effect The resulting brittleness varies with the aging temperature and time Aging at room temperature is very slow, requiring several months to obtain maximum embrittlement As the aging temperature is increased, the time for maximum embrittlement decreases Certain coating treatments, such as hot dip galvanizing, can produce a high degree of embrittlement in areas that were cold worked the critical amount; this can lead to brittle fractures To prevent this problem, the part can be annealed before galvanizing Alternatively, the use of sheet steels containing elements that tie up nitrogen, for example, aluminum, titanium, zirconium, vanadium, or boron, will prevent strain-age embrittlement

Strain aging may also lead to stretcher-strain marks (Lüders bands) on cold-formed low-carbon sheet steel components These marks are cosmetic defects, rather than cracks, but their presence on formed parts is unacceptable (Fig 9) During tensile loading, sheet steel that is susceptible to this defect will exhibit nonuniform yielding followed by uniform elongation The elongation at maximum load and the total elongation are reduced, decreasing cold formability In a nonaluminum-killed sheet steel, a small amount of deformation, typically 1% reduction, will suppress the yield point for several months This process is referred to as roller levelling or temper rolling (Ref 31) This process is more effective in eliminating the sharp yield point and preventing strain aging than stretching the steel through the Lüders strain, which requires about 4 to 6% reduction However, if the material is not formed within the safe period, discontinuous yielding will eventually return and impair formability

Fig 9 Example of stretcher-strain marks (Lüders bands) on a cold-formed steel part

Results of one study illustrate the influence of strain aging on mechanical properties (Ref 32) Three steels made by different processes were evaluated: Steel A, silicon and aluminum deoxidized steel; Steel B, capped open hearth steel; and Steel C, capped Bessemer steel Steel C had the highest nitrogen content Steels B and C had low aluminum contents, while steel A had sufficient aluminum to tie up the nitrogen Strips of each were normalized and loaded in tension to a permanent strain of 10% The strips were held at 25, 230, 480, and 650 °C (75, 450, 900, and 1200 °F) for various lengths

of time (≤25,000 h at 25 and 230 °C, or 75 and 450 °F; ≤10 h at 480 °C, or 900 °F; and 2 h at 650 °C, or 1200 °F) Hardness, tensile properties, and impact properties (half-width Charpy V-notch specimens) were determined at different aging times

Figure 10 shows the impact test results for steels A, B, and C strained 10% in tension and aged at room temperature up to 25,000 h The impact curves are shifted with aging at room temperature for all three steels; steel A exhibits the best aged toughness, and steel C the poorest Figure 11 shows the increase in hardness for steels A, B, and C aged for times up to

25,000 h at 25 °C (75 °F) and 230 °C (450 °F) Room-temperature aging produced a gradual increase in hardness with time The maximum hardness was about the same and was reached quickest by steel C and slowest by steel A The hardness increase with aging at 230 °C (450 °F) was constant for steel A and slowly decreased with aging for steels B and

C

Fig 10 Influence of straining in tension and aging at 24 °C (75 °F) on the Charpy V-notch (half width) impact

strength for three steels (a) Steel A, silicon and aluminum killed, 0.25% C with 0.013% Al and 0.011% N (b) Steel B, capped open hearth steel, 0.07% C with 0.005% Al and 0.005% N (c) Steel C, capped Bessemer steel, 0.08% C with 0.006% Al and 0.016% N All three steels were strained 10% and aged Source: Ref 32

Fig 11 Increase in hardness for steels A, B, and C from Fig 10 after straining in tension (10%) and aging at

24 and 230 °C (75 and 450 °F) for up to 25,000 h Source: Ref 32

In low-carbon steels, strain aging is caused chiefly by the presence of interstitial solutes (carbon and nitrogen), although hydrogen is known to produce a lesser effect at low temperatures These interstitial solutes have high diffusion coefficients in iron and high interaction energies with dislocations The change in mechanical properties of low-carbon rimming steels with different carbon and nitrogen contents that were prestrained 4% and aged various times at 60 °C (140

°F) has been demonstrated (Fig 12) (Ref 34) This work clearly demonstrates the detrimental influence of higher carbon and nitrogen contents on strain aging The solubilities of carbon and nitrogen in iron are quite different Nitrogen solubility is high in the temperature range where rapid precipitation occurs; the solubility of carbon, in equilibrium with cementite, is very low Therefore, strain aging that is due to carbon at temperatures below 100 °C (210 °F) is insignificant However, above 100 °C (210 °F), ε carbide can redissolve and produce substantial strain aging (Ref 35) Strain aging attributable to nitrogen is caused by nitrogen that is not tied up with strong nitride formers, for example, aluminum, titanium, zirconium, vanadium, or boron

Fig 12 Influence of grain size and aging time at 60 °C (140 °F) on the mechanical property changes caused by

strain aging A 0.038C-0.0042N-0.001Al steel was quenched from 200 °C (390 °F), prestrained 4%, and tested after different aging times Grain sizes, in grains/mm 2 (ASTM number), were: specimen 1, 50 (2.7); specimen

2, 195 (4.7); specimen 5, 1850 (7.9) Source: Ref 34

Dynamic Strain Aging. Strain aging can also occur dynamically, that is, aging occurs simultaneously with straining

In this case, the effective strain rate, that is, the dislocation velocity, controls the extent of aging of a particular steel For normal tensile strain rates, dynamic strain aging occurs in the temperature range of 100 to 300 °C (210 to 570 °F) (which includes temperatures at which blue brittleness occurs) If the interstitial solute content is substantial, dynamic strain aging may be observed at room temperature At very high strain rates, as in impact testing, dynamic strain aging is observed at temperatures above 400 °C (750 °F), up to about 670 °C (1240 °F) Reference 33 presents dynamic strain aging results for five carbon steels Carbon and nitrogen are, again, the most important elements in dynamic strain aging Nitrogen is more important than carbon because of the lower solubility of carbon

Aluminum Nitride Embrittlement

It is well known that aluminum nitride precipitation in aluminum-killed steels can cause embrittlement and fracture Several types of problems due to aluminum nitride precipitation have been observed: intergranular fractures in castings (Ref 36, 37, 38, 39, 40, 41, 42, 43), panel cracking in ingots (Ref 44, 45, 46, 47), and reduced hot ductility (Ref 48, 49,

50, 51, 52, 53)

Intergranular fracture in castings, in both the as-cast and heat-treated conditions, have been sporadically observed for many years The fractures occur at the primary austenite grain boundaries formed during solidification In as-cast specimens, ferrite films are generally observed at these grain boundaries The incidence of such cracking has been shown

to increase with increases in aluminum and nitrogen contents and with slower cooling rates after casting

It has been claimed that additions of titanium, zirconium, boron, sulfur, molybdenum, or copper decrease the tendency for cracking (Ref 36) The cooling rate between 1150 and 700 °C (2100 and 1290 °F) is important in controlling the amount

of aluminum nitride precipitation The minimum amount of aluminum nitride required to produce intergranular fracture is lower for alloy steels than for plain carbon steels (0.002 versus 0.004%) (Ref 37) Minimizing the nitrogen content, using the lowest possible amount of aluminum for deoxidation, and increasing the cooling rate after solidification are recommended, and it has been demonstrated that cracking can be prevented by deoxidizing with titanium or zirconium or

by combined titanium and aluminum (Ref 37)

Some researchers have claimed that higher levels of phosphorus and sulfur reduce the susceptibility to aluminum nitride intergranular fractures (Ref 38) Nitrogen content has been shown to be more important than aluminum content because aluminum is always present in amounts greater than that required to tie up all of the nitrogen (Ref 41) Higher aluminum contents do, however, increase the solubility temperature of aluminum nitride and provide an additional driving force for precipitation

Aluminum nitride is known to precipitate with one of two morphologies: plates or dendritic arrays (Ref 38) The dendritic form of aluminum nitride found on the intergranular fracture surfaces of aluminum nitride embrittled castings precipitates from the liquid near the conclusion of solidification (Ref 42) These aluminum nitride dendrites may be nucleation sites for platelike aluminum nitride that precipitates after solidification The platelike aluminum nitride produces that small, shiny fracture surface facets that are generally observed (Ref 43)

Panel Cracking. Panel cracks are longitudinally oriented surface cracks on the side face of an ingot that generally form near the center of the face and can extend to the midradius of the ingot (Ref 44) Such cracks can occur on more than one ingot face and can run the entire length of the ingot Panel cracks are observed in aluminum-killed steel ingots, particularly those with 0.4 to 0.7% C (plain carbon steels) or those with somewhat lower carbon contents for alloy steel ingots These carbon contents lead to ferrite grain-boundary network films with predominantly pearlitic matrix structures The susceptibility to panel cracking varies with melt practice: electric arc furnace steels are most prone; basic open hearth, basic oxygen furnace, and acid open hearth steels are less prone Large ingots are more susceptible than small ingots Stripping of the ingot at as hot a temperature as possible reduces susceptibility Again, aluminum and nitrogen contents are very important Panel cracking is not observed with less than 0.015% Al but occurs with increasing frequency with increasing aluminum content above this level Also, for a given aluminum content, increasing the nitrogen above 0.005% increases panel cracking Crack surfaces are oxidized but not decarburized Because the intergranular cracks generally propagate along ferrite/pearlite interfaces, it has been suggested that cracking occurs at relatively low temperatures, probably below 850 °C (1560 °F) (Ref 44)

A statistical study of panel cracking found that the only significant variable was the level of the aluminum addition (Ref 45) It was suggested that cracking began internally and propagated to the surface because of cooling-induced stresses In

a study of panel cracking in two alloy steels, one containing 0.025% Al and 0.008% N and the other having 0.004% Al

and 0.006% N, the former exhibited panel cracking, and the latter did not Additions of aluminum and titanium to the crack-prone steel composition resulted in the elimination of grain-boundary ferrite networks and freedom from cracking (Ref 46)

A statistical study of panel cracking in low-carbon plate steels found that the steels did not exhibit grain-boundary ferrite networks, but rolled bloom surfaces were heavily cracked in some cases Some of the observations in this study differ from those of other authors who have studied panel cracking However, the statistical analysis showed that cracking increased with aluminum content and track time (time between ingot stripping and charging into the soaking pit) Cracking was minimized by holding the aluminum contents to 0.030% or less (Ref 47)

Reduced Hot Ductility. Numerous studies have demonstrated that increasing aluminum and nitrogen contents degrade hot ductility; this influence is most pronounced in the temperature range where aluminum nitride precipitation is greatest (Ref 48, 49, 50, 51, 52, 53) One author found that 3.4% Ni raised the solution temperature of aluminum nitride in En 36 alloy steel by about 100 °C (210 °F) compared to plain carbon steels with the same aluminum and nitrogen levels (Ref 48) Another study evaluated the hot ductility of chromium-molybdenum-vanadium turbine rotor steels with 0.002 to 0.066% Al and 0.007 to 0.014% N This study showed that nitrogen and residuals (copper, tin, antimony, and arsenic) reduced the hot ductility The addition of titanium and/or boron improved the hot ductility in the test range (800 to 1000

°C, or 1470 to 1830 °F) where high nitrogen contents were detrimental (Ref 49)

In hot ductility tests on carbon-manganese steels containing 0.032 to 0.073% Al and 0.0073 to 0.0105% N, aluminum nitride was found to reduce deformability and increase the resistance to the deformation of austenite These trends were enhanced as the aluminum nitride particle size decreased (Ref 50) Other hot ductility tests on low-carbon steels also showed a large decrease in hot ductility because of aluminum nitride precipitation, depending on the volume fraction and size of the precipitates (Ref 51) Five steels were tested containing different levels of soluble aluminum (aluminum not tied up with oxygen) and with nitrogen present as aluminum nitride Figure 13 shows the reduction in area for hot tensile tests over a range of temperatures for five steels Steels A and B showed a large decrease in percent reduction in area with increasing temperature; however, steel C did not exhibit such embrittlement The difference between steels A, B, and C, which all had high levels of nitrogen in the form of aluminum nitride, was the particle size The smaller aluminum nitride particles in steels A and B were detrimental to hot ductility, while the larger particles in steel C did not cause embrittlement; this agrees with the findings in Ref 50 Steel E contained 0.06% Ti and exhibited the best hot ductility

Composition Particle size

(a) As AIN

(b) No data

Fig 13 Elevated-temperature tensile test results for five plain carbon steels containing various amounts of

aluminum nitride The nitrogen content (in ppm) of the steels in the form of aluminum nitride was: A, 80; B, 70; C, 72; D, 2; E, 1 Source: Ref 51

An evaluation of the hot ductility of low-carbon killed steels found a substantial reduction in fracture strain between 700 and 800 °C (1290 and 1470 °F) This reduction was most pronounced between 750 and 775 °C (1380 and 1425 °F), where ferrite formed along the prior-austenite grain boundaries The intergranular fracture surfaces exhibited aluminum nitride and manganese sulfide precipitates (Ref 53)

Hot ductility is also impaired by high residual impurity contents, chiefly those of copper and tin (Ref 54, 55, 56, 57, 58,

59, 60) Examinations of forging and rolling defects have frequently revealed concentrations of elemental copper, high levels of copper and tin, or copper in the scale Tin residual levels are normally much lower than copper residual levels, but there appears to be a synergistic effect between copper and tin that enhances embrittlement (Ref 54) An examination

of longitudinal cracks in medium- and high-carbon steels found that cracking occurred between 700 and 500 °C (1290 and 930 °F) and depended on high copper and tin contents Aluminum was not added to these steels, but grain-boundary ferrite networks were present The copper and tin were segregated to the ferrite networks (Ref 55)

Nickel, copper, tin, antimony, and arsenic often become enriched in the subscale layer at the surface of steels heated for forging and rolling in oxidizing atmospheres It has been shown that tin, antimony, and arsenic residuals alter the solubility of copper in austenite during high-temperature soaking Because the hot-working temperature is usually above the melting point of elemental copper, liquid copper is produced that will penetrate the austenite grain boundaries and cause cracking by liquid metal embrittlement Nickel and molybdenum concentrate with copper and raise the melting point of copper; tin, antimony, and arsenic also concentrate at the scale/metal interface and lower the melting point of copper If copper is not present, tin, antimony, and arsenic have little detrimental effect on hot workability (Ref 56)

It has been shown that tin reduces the solubility of copper in austenite, which is probably more important than its influence on the melting point of copper (Ref 56) Tin reduces the solubility of copper in austenite by a factor of three; therefore, when tin is present, molten copper can form at the surface at lower bulk copper contents Nickel reduces copper-induced hot shortness, manganese slightly increases hot shortness, arsenic slightly more detrimental than manganese, and tin and antimony are extremely detrimental to copper-induced hot shortness (Ref 58) Chromium decreases the solubility of copper in austenite and increases the susceptibility to copper-induced hot shortness, although its influence is small (Ref 59)

Graphitization

In the early 1940s, several failures of welded joints in high-pressure steam lines occurred because of graphite formation in the region of the weld heat-affected zone that had been heated during welding to the critical temperature of the steel used (Ref 61, 62, 63, 64, 65) Extensive surveys of carbon and carbon-molybdenum steel specimens removed from various types of petroleum refining equipment revealed graphite in about one-third of the 554 specimens examined (Ref 61, 64)

In most cases, graphite formation did not occur until about 40,000 h or more at temperatures from 455 to 595 °C (850 to

1100 °F) Aluminum-killed carbon steels were susceptible, but silicon-killed or low-aluminum killed steels were not The C-0.5Mo steels were more resistant to graphitization than carbon steels, but they were similarly influenced by the nature

of the deoxidation practice Chromium additions and stress relieving at 650 °C (1200 °F) both retarded graphitization

An examination of the graphitization susceptibility of a number of alloy steels showed that chromium-molybdenum steels used for steam piping, either annealed or normalized, were resistant to graphitization Nickel and nickel-molybdenum steels did graphite during high-temperature exposure Chromium-bearing steels did not graphitize and appeared to be quite stable (Ref 65)

The deoxidation practice used in making the steels is the most important parameter influencing graphitization High levels

of aluminum deoxidation strongly promote graphitization for both carbon and carbon-molybdenum steels used for steam lines While nitrogen appears to retard graphitization, high levels of aluminum remove nitrogen and thus promote graphitization (Ref 66) Molybdenum additions (0.5%) help stabilize cementite but do not fully offset the influence of high aluminum additions Manganese and silicon both affect graphitization, but their influence is small at the levels used

in these alloys Chromium appears to be the best alloy addition for stabilizing carbides

Tensile tests of affected steam piping indicate that the graphite present did not affect the tensile strength Charpy V-notch impact strength, however, was reduced substantially Localized graphitization near a welded joint appears to be much more damaging to pipe behavior than general, uniform graphitization The localized graphitization apparently produces notches that concentrate stress and reduce load-bearing capability

References cited in this section

19 R.L Kenyon and R.S Burns, Testing Sheets for Blue Brittleness and Stability Against Changes Due to

Aging, Proc ASTM, Vol 34, 1934, p 48-58

20 E.O Hall, The Deformation of Low-Carbon Steel in the Blue-Brittle Range, J Iron Steel Inst., Vol 170,

April 1952, p 331-336

21 G Mima and F Inoko, A Study of the Blue-Brittle Behavior of a Mild Steel in Torsional Deformation,

Trans JIM, Vol 10, May 1969, p 227-231

22 T Takeyama and H Takahashi, Strength and Dislocation Structures of α-Irons Deformed in the

Blue-Brittleness Temperature Range, Trans ISIJ, Vol 13, 1973, p 293-302

23 A.L Tsou et al., The Quench-Aging of Iron, J Iron Steel Inst., Vol 172, Oct 1952, p 163-171

24 T.C Lindley and C.E Richards, The Effect of Quench-Aging on the Cleavage Fracture of a Low-Carbon

Steel, Met Sci J., Vol 4, May 1970, p 81-84

25 A.S Keh and W.C Leslie, Recent Observations on Quench-Aging and Strain Aging of Iron and Steel, in

Materials Science Research, Vol 1, Plenum Publishing, 1963, p 208-250

26 E.R Morgan and J.F Enrietto, Aging in Steels, in AISI 1963 Regional Technical Meeting, American Iron

and Steel Institute, 1964, p 227-252

27 E Stolfe and W Heller, The State of Knowledge of the Aging of Steels: I, Fundamental Principles, Stahl und Eisen, Vol 90 (No 16), 1970, p 861-868

28 G Lagerberg and B.S Lement, Morphological and Phase Changes During Quench-Aging of Ferrite

Containing Carbon and Nitrogen, Trans ASM, Vol 50, 1958, p 141-162

29 J.D Baird, Strain Aging of Steel A Critical Review, Iron Steel, Vol 36, 1963, p 186-192, 326-334,

368-374, 400-405, and 450-457

30 J.D Baird, The Effects of Strain-Aging Due to Interstitial Solutes on the Mechanical Properties of Metals,

Met Rev., Vol 16, Feb 1971, p 1-18

31 R.D Butler and D.V Wilson, The Mechanical Behavior of Temper Rolled Steel Sheets, J Iron Steel Inst.,

Vol 201, Jan 1963, p 16-33

32 F Garofalo and G.V Smith, The Effect of Time and Temperature on Various Mechanical Properties During

Strain Aging of Normalized Low Carbon Steels, Trans ASM, Vol 47, 1955, p 957-983

33 C.C Li and W.C Leslie, Effects of Dynamic Strain Aging on the Subsequent Mechanical Properties of

Carbon Steels, Metall Trans A, Vol 9A, Dec 1978, p 1765-1775

34 D.V Wilson and B Russell, The Contribution of Precipitation to Strain Aging in Low Carbon Steels, Acta Metall., Vol 8, July 1960, p 468-479

35 E.T Stephenson and M Cohen, The Effect of Prestraining and Retempering on AISI Type 4340, Trans ASM, Vol 54, 1961, p 72-83

36 C.H Lorig and A.R Elsea, Occurrence of Intergranular Fracture in Cast Steels, Trans AFS, Vol 55, 1947, p

160-174

37 B.C Woodfine and A.G Quarrell, Effects of Al and N on the Occurrence of Intergranular Fracture in Steel

Castings, J Iron Steel Inst., Vol 195, Aug 1960, p 409-414

38 J.A Wright and A.G Quarrell, Effect of Chemical Composition on the Occurrence of Intergranular

Fracture in Plain Carbon Steel Castings Containing Aluminum and Nitrogen, J Iron Steel Inst., Vol 200,

41 N.H Croft, Use of Solubility Data to Predict the Effects of Aluminum and Nitrogen Contents on the

Susceptibility of Steel Castings to Intergranular Embrittlement, Met Technol.,, Vol 10, Aug 1983, p

44 S.C Desai, Longitudinal Panel Cracking in Ingots, J Iron Steel Inst., Vol 191, March 1959, p 250-256

45 E Colombo and B Cesari, The Study of the Influence of Al and N on the Susceptibility to Crack Formation

of Medium Carbon Steel Ingots, Metall Ital., Vol 59, 1967, p 71-75

46 L Ericson, Cracking in Low Alloy Aluminum Grain Refined Steels, Scand J Metall., Vol 6, 1977, p

116-124

47 R Sussman et al., Occurrence and Control of Panel Cracking in Aluminum Containing Steel Heats, in Mechanical Work and Steel Processing, Vol 17, American Institute of Mining, Metallurgical, and

Petroleum Engineers, 1979, p 49-78

48 L.A Erasmus, Effect of Aluminum Additions on Forgeability, Austenite Grain Coarsening Temperature,

and Impact Properties, J Iron Steel Inst., Vol 202, Jan 1964, p 32-41

49 R Harris and L Barnard, Experiences of Hot Shortness in the Forging of Certain Low-Alloy Steels, in

Deformation Under Hot Working Conditions, SR 108, Iron and Steel Institute, 1968, p 167-177

50 F Vodopivec, Influence of Precipitation and Precipitates of Aluminum Nitride on Torsional Deformability

of Low-Carbon Steel, Met Technol., Vol 5, April 1978, p 118-121

51 G.D Funnell and R.J Davies, Effect of Aluminum Nitride Particles on Hot Ductility of Steel, Met Technol., Vol 5, May 1978, p 150 153

52 G.D Funnell, Observations on Effect of Aluminum Nitride on Hot Ductility of Steel, in Hot Working and Forming Processes, Book 264, The Metals Society, 1980, p 104-107

53 K Yamanaka et al., Relation Between Hot Ductility and Grain-Boundary Embrittlement of Low-Carbon Killed Steels, Trans ISIJ, Vol 20, 1980, p 810-816

54 S.L Gertsman and H.P Tardif, Tin and Copper in Steel: Both Are Bad, Together They're Worse, Iron Age,

Vol 169 (No 7), Feb 14, 1952, p 136-140

55 P Bjornson and H Nathorst, A Special Type of Ingot Cracks Caused by Certain Impurities, Jernkontorets Ann., Vol 139 (No 6), 1955, p 412-438

56 D.A Melford, Surface Hot Shortness in Mild Steel, J Iron Steel Inst., Vol 200, April 1962, p 290-299

57 A Nicholson and J.D Murray, Surface Hot Shortness in Low-Carbon Steel, J Iron Steel Inst., Vol 203, Oct

60 W.J.M Salter, Effect of Mutual Additions of Tin and Nickel on the Solubility and Surface Energy of

Copper in Mild Steel, J Iron Steel Inst., Vol 207, Dec 1969, p 1619-1623

61 H.J Kerr and F Eberle, Graphitization of Low-Carbon and Low-Carbon-Molybdenum Steels, Trans

ASME, Vol 67, 1945, p 1-46

62 S.L Hoyt et al., Summary Report on the Joint E.E.I.-A.E.I.C Investigation of Graphitization in Piping, Trans ASME, Vol 68, Aug 1946, p 571-580

63 R.W Emerson and M Morrow, Further Observations of Graphitization in Aluminum-Killed

Carbon-Molybdenum Steel Steam Piping, Trans AIME, Vol 68, Aug 1946, p 597-607

64 J.G Wilson, Graphitization of Steel in Petroleum Refining Equipment, Weld Res Counc Bull., No 32, Jan

1957, p 1-10

65 A.B Wilder et al., Stability of AISI Alloy Steels, Trans AIME, Vol 209, Oct 1957, p 1176-1181

66 E.J Dulis and G.V Smith, Roles of Aluminum and Nitrogen in Graphitization, Trans ASM, Vol 46, 1954,

p 1318-1330

Embrittlement of Steels

George F Vander Voort, Carpenter Technology Corporation

Overheating

The overheating of steels occurs when they are heated to excessively high temperatures prior to hot working (Ref 67, 68,

69, 70, 71, 72, 73, 74, 75, 76) Heating at even higher temperatures will cause incipient grain-boundary melting, a problem known as burning Thus, overheating occurs in the temperature range between the safe range normally used prior

to hot working and the higher range where liquation begins Hot working after burning generally results in the tearing or rupture of the steel due to the liquid in the grain boundaries Hot working after overheating generally does not result in cracking; if sufficient hot reduction occurs, the influence of overheating may be minor or negligible If the degree of hot reduction is small, mechanical properties, chiefly toughness and ductility, will be affected

Fracture surfaces of overheated steels given limited hot reduction often exhibit a coarse-grain faceted appearance Such features are most evident after quenching and tempering to develop optimum toughness Other problems, such as aluminum nitride embrittlement, may also produce a faceted fracture surface Additional tests are needed to distinguish among these problems (Ref 76)

Although the mechanical properties of burnt steels are always very poor, the mechanical properties of overheated steels show considerable scatter For tensile tests, the elongation and reduction of area are most affected by overheating and decrease with increasing heating temperature Fracture faceting and substantial decreases in tensile ductility normally are observed after severe overheating

Impact properties are usually more sensitive to overheating than is tensile ductility In examining impact test results, several interrelated features should be examined:

• Change in upper-shelf energy

• Impact strength transition temperature

• Presence of facets

Upper-shelf energy appears to be particularly sensitive to overheating Figure 14 shows impact energy curves for En

111 alloy steel specimens heated to a variety of temperatures from the norma soaking range to the burning range; no hot working was performed The specimens were first heated to 950 °C (1740 °F) for 10 min, transferred to the high-temperature zone for 7 min, then transferred back to the furnace at 950 °C (1740 °F) and held for 50 min They were oil quenched, tempered 1 h at 675 °C (1245 °F), and water quenched to minimize temper embrittlement (more information about temper embrittlement is given in the section "Temper Embrittlement in Alloy Steels" in this article) The use of temperatures up to 1200 °C (2190 °F) produced no change in impact energies, but temperatures above 1200 °C (2190 °F) produced a decrease in upper-shelf energy The burnt specimens displayed a substantial loss in toughness Because the pieces were not forged after the burning treatment, they did not fracture

Fig 14 Impact energy values versus test temperature for En 111 alloy steel specimens heated to the indicated

temperature for 1 h, oil quenched, and tempered for 1 h at 675 °C (1245 °F) Source Ref 68

Presence of Facets. Many of the studies of overheating have been conducted without subsequent hot working after overheating and thus do not simulate actual commercial experience These studies are of limited value Numerous theories have been proposed to explain overheating The examination of facets on fractures of overheated steels reveals that the facet surfaces are covered with fine ductile dimples, and small manganese sulfides can be found within the dimples (because two fracture surfaces are formed, a manganese sulfide particle will be found in either half of the mating dimples after fracture) The facet surfaces correspond to prior-austenite grain surfaces formed during overheating

As the soaking temperature is raised, manganese sulfide in the steel dissolves (that is, the sulfur goes into solution in the austenite) Dissolved sulfur diffuses toward the austenite grain boundaries, where it reprecipitates Therefore, overheating changes the size and distribution of sulfides in the steel The cooling rate through the overheating range also affects the size and dispersion of the intergranular sulfides In commercial practice, the size of the overheated piece, and any externally applied coolant during hot working, will control this cooling rate

Steels containing less than 10 ppm S do not overheat, regardless of the heating temperature up to burning However, this level of sulfur is difficult to obtain Steels with relatively low sulfur contents, (for example, in the range of 0.001 to 0.005%) are being produced in greater quantities today because of the detrimental influence of sulfur on properties However, it has been demonstrated that the minimum temperature at which overheating occurs in low-sulfur steels is lower than that for high-sulfur steels (>0.10% S) Overheating problems thus have been experienced in low-sulfur steels heated at temperatures that usually do not cause overheating (Ref 73, 74, 75) Additions of rare earth elements have been shown to prevent overheating by modifying the solubility of the sulfide formed High-sulfur steels appear to require a greater degree of overheating to cause fracture faceting and impaired properties than do low-sulfur steels

The problem of overheating is complex Sulfide dissolution and reprecipitation at the prior-austenite grain boundaries causes fracture faceting and impairment of properties However, when faceting is observed, mechanical properties may not be significantly affected Overheating and its influence on properties depend on the sulfur content, the soaking temperature, grain size, the cooling rate through the overheating range, and the degree of hot reduction Furthermore, the amount of faceting observed on the test fracture depends on the heat treatment (particularly the tempering temperature), the test temperature, the test specimen orientation, and the amount of deformation after sulfide reprecipitation

References cited in this section

67 A Preece et al., The Overheating and Burning of Steel, J Iron Steel Inst., Part I, Vol 153, 1946, p

237p-254p; and, Part III, Vol 164, 1950, p 37-45

68 I.S Brammar, A New Examination of the Phenomena of Overheating and Burning of Steels, J Iron Steel Inst., Vol 201, Sept 1963, p 752-761

69 R.D.N Lester, Overheating in Steels, Steel Times, Vol 193 (No 513), 15 July 1966, p 96-102

70 G.D Joy and J Nutting, Influence of Intermetallic Phases and Non-Metallic Inclusions Upon the Ductility

and Fracture Behavior of Some Alloy Steels, in Effect of Second-Phase Particles on the Mechanical Properties of Steel, Iron and Steel Institute, 1971, p 95-100

71 T.J Baker and R Johnson, Overheating and Fracture Toughness, J Iron Steel Inst., Vol 211, Nov 1973, p

74 R.C Andrew and G.M Weston, The Effect of Overheating on the Toughness of Low Sulphur ESR Steels,

J Aust Inst Met., Vol 22, Sept-Dec 1972, p 171-176

75 R.C Andrew and G.M Weston, The Effect of the Interaction Between Overheating and Tempering

Temperature on the Notch Toughness of Two Low Sulphur Steels, J Aust Inst Met., Vol 22, Sept-Dec

1972, p 200-204

76 G.E Hole and J Nutting, Overheating of Low-Alloy Steels, Int Met Rev., Vol 29, 1984, p 273-298

Embrittlement of Steels

George F Vander Voort, Carpenter Technology Corporation

Thermal Embrittlement of Maraging Steels

Maraging steels will fracture intergranularly at low impact energies if improperly processed after hot working This problem, known as thermal embrittlement, occurs when maraging steels that have been heated above 1095 °C (2000 °F) are slowly cooled through, or held within, the temperature range of 980 to 815 °C (1800 to 1500 °F) (Ref 77, 78, 79, 80, 81) The embrittlement is caused by the precipitation of TiC and/or Ti(C,N) on the austenite grain boundaries during cooling through, or holding within, the critical temperature range The degree of embrittlement increases with time within the critical range Increased levels of carbon and nitrogen render maraging steels more susceptible to thermal embrittlement Auger analysis has shown that embrittlement begins with the diffusion of titanium, carbon, and nitrogen to the grain boundaries, and observation of TiC or Ti(C,N) precipitates represents an advanced stage of embrittlement

Results of a study on thermal embrittlement demonstrate its influence on an 18Ni(250) maraging steel (Ref 77) Plates 12.7 mm (0.5 in.) thick were rolled with finishing temperatures in the range of 1080 to 870 °C (1980 to 1600 °F) and then cooled by three different methods (water, air, and vermiculite) These results showed that the finishing temperature and cooling rate from the finishing temperature had minor effects on the tensile properties but a significant effect on fracture toughness Figure 15 shows the plane-strain fracture toughness results as a function of finishing temperature (temperature

at the end of rolling) and cooling rate for hot-rolled and aged material (Fig 15a) and for annealed and aged material (Fig

15b and 15c) In general, vermiculite cooling (slow) and a high finishing temperature produced the lowest toughness, except for specimens water quenched after rolling (Fig 15b and 15c) Detailed information on maraging steels is available in the article "Maraging Steels" in this Volume

Fig 15 Influence of mill finishing temperature and manner of cooling on the plane-strain fracture toughness

(KIc ) of 18Ni(250) maraging steel heat treated three ways (a) Hot rolled and aged (b) Annealed at 870 °C (1600 °F) and aged (c) Annealed at 815 °C (1500 °F) and aged Source: Ref 77

References cited in this section

77 G.J Spaeder et al., The Effect of Hot Rolling Variables on the Fracture Toughness of 18Ni Maraging Steel, Trans ASM, Vol 60, 1967, p 418-425

78 D Kalish and H.J Rack, Thermal Embrittlement of 18Ni(350) Maraging Steel, Metall Trans., Vol 2, Sept

1971, p 2665-2672

79 W.C Johnson and D.F Stein, A Study of Grain Boundary Segregants in Thermally Embrittled Maraging

Steel, Metall Trans., Vol 5, March 1974, p 549-554

80 E Nes and G Thomas, Precipitation of TiC in Thermally Embrittled Maraging Steels, Metall Trans A, Vol

Numerous factors can contribute to cracking susceptibility (Ref 82, 83, 84, 85, 86, 87):

• Heat treatment practice

As the carbon content is raised, the Ms and Mf (temperature at which martensite formation ends) temperatures decrease, and the volumetric expansion and transformation stresses accompanying martensite formation increase In general, steels with less than 0.35% C are free of quench cracking problems Such low-carbon steels have higher Ms and Mftemperatures, which allow some stress relief to occur during the quench Also, transformation stresses are lower, and the lower strength of the martensite formed (low-carbon lath martensite) can accommodate the strains more readily than can a higher-carbon steel

Alloy steels with ideal critical diameters of 4 or greater are more susceptible to quench cracking than are hardenability steels Quench crack sensitivity also increases as the severity of the quench rate increases Control of the austenitizing temperature is also important, particularly for high-carbon tool steels Excessive retained austenite and coarse-grain structures promote quench cracking Quench uniformity is important, particularly when liquid quenchants are employed When high-carbon steels are quenched to form martensite, they are in a highly stressed condition Therefore, tempering must be done immediately after quenching to relieve these stresses and minimize the risk of cracking Surface quality is also very important because seams, laps, tool marks, stamp marks, and so on, act as stress concentrators to locate and enhance quench cracking susceptibility (Ref 82)

lower-Quench cracking is a problem that often defies prediction and can be difficult to diagnose Heat treaters have experienced short time periods in which cracking problems occur frequently and then stop for no apparent reason Evidence also indicates that quench cracking can be more frequent for certain heats of steel, again for no obvious reason Instances have

also been documented (Ref 86) in which extensive quench cracking has occurred in material processed from the bottom portion of ingots

Quench crack fractures are always intergranular In quenched and tempered steels, proof of quench cracking is often obtained by opening a crack (if necessary) and visually looking for the temper color typical for the temperature used (Ref

88, 89) If the crack occurs during or after quenching but before tempering, and if the crack is open to the furnace atmosphere, a thin oxide layer will form on the surface The color of the oxide layer depends on its thickness, which in turn depends on the tempering temperature and the steel composition Quench cracks begin at the surface and propagate inward They are usually oriented longitudinally or radially unless located by a change in section size, by surface imperfections, or by changes in surface microstructure (such as an interface between hardened and nonhardened areas)

References cited in this section

82 G.F Vander Voort, Failures of Tools and Dies, in Failure Analysis and Prevention, Vol 11, 9th ed., Metals Handbook, American Society for Metals, 1986, p 563-585

83 L.D Jaffee and J.R Hollomon, Hardenability and Quench Cracking, Trans AIME, Vol 167, 1946, p

s, Met Prog., Vol 65, May 1954, p 113-121

87 T Kunitake and S Sugisawa, The Quench-Cracking Susceptibility of Steel, Sumitomo Search, No 5, May

1971, p 16-25

88 P Gordon, The Temper Colors on Steel, J Heat Treat., Vol 1, June 1979, p 93

89 G.F Vander Voort, Visual Examination and Light Microscopy, in Fractography, Vol 12, 9th ed., Metals Handbook, ASM INTERNATIONAL, 1987, p 91-165

Embrittlement of Steels

George F Vander Voort, Carpenter Technology Corporation

Temper Embrittlement in Alloy Steels

Temper embrittlement also known as temper brittleness*, two-step temper embrittlement, or reversible temper embrittlement is a problem associated with tempered alloy steels that are heated within, or slowly cooled through, a critical temperature range, generally 300 to 600 °C (570 to 1110 °F) for low-alloy steels This treatment causes a decrease

in toughness as determined with Charpy V-notch impact specimens (Ref 90, 91, 92, 93, 94, 95, 96, 97, 98, 99, 100, 101,

102, 103, 104, 105, 106, 107, 108, 109, 110, 111, 112, 113, 114) It is a particular problem for heavy-section components, such as pressure vessels and turbine rotors, that are slowly cooled through the embrittling range after tempering and also experience service at temperatures within the critical range

Temper embrittled steels exhibit an increase in their DBTT and a change in fracture mode in the brittle test temperature range from cleavage to intergranular The DBTT can be assessed in several ways, the most common being the temperature for 50% ductile and 50% brittle fracture (50% fracture appearance transition temperature, or FATT), or the lowest temperature at which the fracture is 100% ductile (100% fibrous criterion) Transition temperatures based on absorbed energy values are not normally employed Temper embrittlement is reversible; that is, the toughness of embrittled steels can be restored by tempering them above the critical region followed by rapid cooling, for example, water quenching This decreases the DBTT and changes the low-temperature (that is, below the 50% FATT) intergranular brittle appearance back to the cleavage mode

Temper embrittlement does not occur in plain carbon steels, only in alloy steels Also, the degree of embrittlement varies with alloy steel composition Therefore, the alloying elements present, and their combinations and levels, are important However, certain impurity elements must be present if temper embrittlement is to occur The embrittling impurities are, in decreasing order of influence on a weight percent basis, antimony, phosphorus, tin, and arsenic Of these elements, phosphorus is most commonly present in alloy steels and it has captured the most attention in research studies Manganese and silicon also increase the susceptibility to embrittlement Although alloy steels are ferritic in the tempered condition, fracture below the DBTT occurs along prior-austenite grain boundaries where both alloying elements and impurity elements are concentrated

Effect of Composition on ∆FATT

The proof that antimony, phosphorus, tin and/or arsenic is an essential ingredient(s) for temper embrittlement was obtained in the late 1950s (Ref 90, 91) The change in 50% FATT and 100% fibrous FATT with isothermal aging at 450

°C (230 °F) for up to 1000 h for nickel-chromium and nickel-chromium-molybdenum laboratory heats with controlled compositions and impurity levels was determined Figures 16 and 17 show some test results for the influence of antimony, phosphorus, tin, and arsenic and aging at 450 °C (840 °F) for 1000 h (Ref 90, 91) Embrittlement was greater for the nickel-chromium steels than for the nickel-chromium-molybdenum steels because of the beneficial influence of molybdenum The nickel-chromium steels also showed substantially greater embrittlement from the manganese addition than did the nickel-chromium-molybdenum steels The addition of about 0.7% Si had a smaller and similar embrittling influence for both grades

Fig 16 Influence of phosphorus, antimony, arsenic, and tin impurity elements on the temper embrittlement

susceptibility of nickel-chromium experimental steels based on the change in (a) 50% FATT and (b) 100%

fibrous FATT after aging at 450 °C (840 °F) for 1000 h Source: Ref 91

Fig 17 Influence of phosphorus, antimony, arsenic, and tin impurity elements on the temper embrittlement

susceptibility of nickel-chromium-molybdenum experimental steels based on the change in (a) 50% FATT and (b) 100% fibrous FATT after aging at 450 °C (840 °F) for 1000 h Source: Ref 91

The important role of alloying elements has been clearly demonstrated in tests that also used heats of controlled composition (Ref 92) The tests were performed using 0.4% C alloy steels containing nickel, chromium, or nickel and chromium, as well as a plain carbon steel Controlled additions of antimony, phosphorus, tin, and arsenic were made to these compositions The 50% FATT was evaluated for each composition after heat treatment (870 °C, or 1600 °F, for 1 h, oil quench; 625 °C, or 1155 °F, for 1 h, water quench) and after step-cool embrittlement Figure 18 shows the results for additions of antimony to plain carbon and alloy steels of various analyses The bars show the 50% FATT after tempering

at 625 °C (1155 °F) (left end not embrittled) and after step cooling (right end embrittled); the value under the bar is the shift in 50% FATT The addition of 600 to 800 ppm Sb to the 0.4C-3.5Ni-1.7Cr steel caused a shift in transition temperature of 695 °C (1285 °F) The same steel, but without carbon, exhibited a shift of 222 °C (432 °F), but its hardness was much lower (~80 HRB versus 27 HRC) The steels with only nickel and carbon or chromium and carbon and 600 to 800 ppm Sb exhibited much less embrittlement, while the plain carbon steel was not embrittled by antimony

Fig 18 Influence of alloying elements on the temper embrittlement of steels (compositions given in

accompanying table) containing 600 to 800 ppm Sb The left end of bar gives the nonembrittled DBTT; the right end of bar gives the DBTT after embrittlement (except for line F, which is reversed) Value between bar ends is the shift in 50% FATT Source: Ref 92

Figure 19 shows the results for additions of about 500 ppm each of phosphorus, tin, and antimony to the chromium-carbon, nickel-carbon, and chromium-carbon steels The nonembrittled toughnesses of the nickel-chromium-carbon-phosphorus and chromium-carbon-phosphorus alloys were poorer than those of the other alloys shown, probably because of the segregation of phosphorus in austenite Phosphorus also embrittled the chromium-carbon-phosphorus alloy much more than the nickel-carbon-phosphorus alloy This is due to an interaction between chromium and phosphorus Tin appears to embrittle the nickel-chromium-carbon alloy more than phosphorus in that the change in fracture appearance transition temperature (∆FATT) was greater However, the grain size of the nickel-chromium-carbon-phosphorus alloy was ASTM No 8, while that of the nickel-chromium-carbon-tin alloy was ASTM No 6 Also, the nonembrittled toughness of the alloy containing phosphorus was much poorer The 50% FATT values for these two compositions are nearly identical and would be even closer if the grain sizes were the same Otherwise, it appears that tin embrittled the nickel-carbon alloy more than phosphorus, while phosphorus embrittled the chromium-carbon alloy more than tin Arsenic was a much weaker embrittler The results in Fig 18 and 19 clearly show that the combination of nickel and chromium resulted in much greater embrittlement, particularly for additions of antimony and tin The results for phosphorus show that a strong interaction exists between chromium and phosphorus, while phosphorus causes little embrittlement in nickel steels not containing chromium

Fig 19 Influence of alloying elements on the temper embrittlement of steels (compositions given in

accompanying tables) (a) Steel containing 500 to 600 ppm P (b) Steel containing 460 to 480 ppm Sn (c) Steel containing 500 to 530 ppm As The left end of bar gives the nonembrittled DBTT; the right end of bar gives the DBTT after embrittlement Value between bar ends is the shift in 50% FATT Source: Ref 92

The beneficial influence of molybdenum on phosphorus-induced temper embrittlement has been known for many years (Ref 93, 94) It has also been known that carbon steels are immune to temper embrittlement, but that substantial additions

of manganese cause susceptibility to this problem In addition, high levels of manganese in alloy steels have been known

to render them more susceptible to temper embrittlement

Reference 95 includes an evaluation of the addition of 0.6% Mo, along with controlled antimony, phosphorus, tin, arsenic, and other elements, to AISI 3340 (3.5Ni-1.7Cr-0.4C) The addition of molybdenum eliminated or greatly reduced embrittlement due to step cooling for additions of antimony, tin, and arsenic, but not for additions of phosphorus The addition of 0.7% Mn to this steel produced substantial embrittlement, which was largely eliminated when 0.6% Mo was added In this work, it was shown that phosphorus segregates to the austenite grain boundaries during austenitization; antimony does not do this This work also clearly showed that manganese is an embrittling element, not merely an enhancer of embrittlement

Later work was conducted to clarify these results (Ref 96) The earlier work had employed specimens with a very coarse grain size (Ref 95) Results with somewhat finer grain sizes, although still rather coarse, showed that 0.5 to 0.6% Mo additions would prevent temper embrittlement caused by phosphorus in 3.5Ni-1.7Cr-0.2C steels for aging times up to

1000 h at 475 and 500 °C (885 and 930 °F) The influence of molybdenum on the prevention of temper embrittlement appears to depend on how much of it is dissolved in the matrix as opposed to how much is tied up in carbides As more molybdenum becomes tied up in the carbides, the beneficial influence of molybdenum decreases Therefore, depending

on the temperatures experienced and the presence of other strong carbide formers, molybdenum may or may not be able

to suppress temper embrittlement

A major advancement in understanding temper embrittlement was brought about by the development of Auger electron spectroscopy and its application to embrittlement studies in 1969 (Ref 97, 98, 99) This work permitted the direct chemical analysis of segregants on the intergranular fracture surfaces of embrittled specimens Such work has shown that the embrittling impurity elements are segregated to within the first few monolayers of the embrittled grain boundaries The degree of enrichment of these elements may be 100 to 103 times the bulk concentration Alloy element segregation at these boundaries was also detected However, the concentration of these alloying elements was found to be only 2 to 3 times that of the bulk concentration, and the concentration profile from the grain boundary into the grain interior was much shallower than for the impurity elements Figure 20 shows an example of Auger analysis of antimony, sulfur, and phosphorus segregated to either fracture or free surfaces (Ref 99, 100) These results were obtained by alternate argon ion sputtering (depth profiling) and analysis

Fig 20 Normalized intensities of Auger peaks (as a function of depth below the surface) from antimony, sulfur,

and phosphorus segregated to grain boundaries or free surfaces (depth profiling by argon ion sputtering) Source: Ref 99

The degree of embrittlement also depends upon the time at temperature within the critical region Extensive isothermal embrittlement studies were performed on a heat of AISI/SAE 3140 alloy steel (0.39C-0.79Mn-0.30Si-0.028S-0.015P-1.26Ni-0.77Cr-0.02Mo, ASTM No 8 grain size) (Ref 101) Specimens were austenitized at 900 °C (1650 °F) for 1 h, water quenched, and tempered at 675 °C (1245 °F) for 1 h, and water quenched This resulted in a hardness of 23 HRC and a transition temperature (100% fibrous criterion) of -83 °C (-117 °F) Specimens were aged at temperatures from 325

to 650 °C (615 to 1200 °F) for times ranging from 4 min to 240 h Figure 21 shows the time-temperature-embrittlement diagram developed from this work

Fig 21 Time-temperature diagram for isothermally temper-embrittled AISI/SAE 3140 alloy steel showing

constant embrittlement levels (100% fibrous FATT) for quenched and tempered (675 °C, or 1245 °F, for 1 h), specimens Source: Ref 101

Because of some minor differences in hardness and apparent inconsistencies in the results for aging at 525 °C (975 °F), additional work was performed (Ref 102) Minor corrections were made in the transition temperatures to account for hardness differences (Fig 22) Only limited tests were done at temperatures above 575 °C (1065 °F), but the work performed indicated that temper embrittlement occurred at tempering temperatures up to 675 °C (1245 °F), which is close

to the lower critical temperature The diagrams exhibit the classic C-curve appearance The nose of the curve is at 550 °C (1020 °F), but maximum embrittlement (∆FATT ≈100 °C, or 210 °F) was obtained at 475 to 500 °C (885 to 930 °F) after

240 h aging (the longest time used) Tests done using step cooling showed that the degree of embrittlement that occurred was greater than would have been predicted from the isothermal data

Fig 22 Revised time-temperature diagram for temper-embrittled AISI/SAE 3140 alloy steel Source: Ref 102

The above data was analyzed to predict the degree of grain-boundary phosphorus segregation in this steel (analysis for antimony, tin, and arsenic was not performed, but the amount of these elements was assumed to be very low) (Ref 103) Auger analysis of similar steels was used to calculate the monolayers of phosphorus segregated to the prior-austenite grain boundaries (Fig 23)

Fig 23 Time-temperature diagram for the segregation of phosphorus in temper-embrittled AISI/SAE 3140

alloy steel The numbers next to the curves describe the degree of phosphorus segregated during the embrittlement treatment (not including the 0.06 monolayers of phosphorus segregated prior to the isothermal aging treatments) Source: Ref 103

Microstructure and Grain Size. It is well known that microstructure influences the susceptibility to temper embrittlement and the resulting degree of embrittlement Because the impurity and alloying elements segregate to the

prior-austenite grain boundaries, grain size has an influence As the grain size becomes larger, the grain-boundary surface area decreases Therefore, for a fixed level of impurities and constant embrittling temperatures and time, there will be greater coverage of the grain boundaries in a coarse-grain steel than in a fine-grain steel However, the distance over which the impurities must diffuse increases as the grain size becomes larger Nevertheless, coarse-grain steels are recognized to be more severely embrittled than fine-grain steels Figure 24 shows the results of aging a 0.33C - 0.59Mn - 0.03P - 0.031S - 0.27Si - 2.92Ni - 0.87Cr steel for various times at 500 °C (930 °F) followed by water quenching Prior to aging, the specimens had been austenitized (at 850 °C, or 1560 °F, for fine-grain specimens; at 1200 °C, or 2190 °F, for coarse-grain specimens) and tempered at 650 °C (1200 °F) for 1 h; they were oil quenched after tempering The coarse-grain specimens were embrittled to a greater extent than the fine-grain specimens Similar results have been obtained by others (for example, Ref 105)

Fig 24 Influence of prior-austenite grain size on the temper embrittlement of a nickel-chromium alloy steel

that was heat treated to produce two levels of grain size The alloy was tempered at 650 °C (1200 °F) and aged various times at 500 °C (930 °F) (a) Actual 100% fibrous FATT (b) Change in 100% fibrous FATT Source: Ref

104

Matrix microstructures are also important because they control toughness, for both nonembrittled and embrittled steels Most studies have evaluated temper embrittlement of martensitic specimens, but a few have compared results for a variety

of microstructures as a function of hardness In general, tempered martensite is more susceptible than tempered bainite to temper embrittlement, but tempered bainite is more susceptible than pearlitic-ferritic structures This analysis is somewhat misleading, however, because nonembrittled tempered martensite is much tougher than nonembrittled bainite at the same hardness, and after embrittlement, tempered martensite is still tougher than tempered bainite The same holds true when comparing bainitic and pearlitic-ferritic microstructures This has been demonstrated for chromium-molybdenum-vanadium steels, as shown in Fig 25 (Ref 106) The alloy composition was 0.3C-0.91Mn-0.27Si-0.15Ni-1.3Cr-1.2Mo-0.31V-0.025P-0.0045S-0.005 As-0.0008Sb-0.027Sn The toughness of embrittled tempered martensite was better than that of nonembrittled tempered bainite over the hardness range evaluated Also, the shift in 50% FATT for tempered martensite increase with hardens and was greater than that for tempered bainite Only one hardness level was obtained for the ferrite-pearlite condition, and the shift in 50% FATT because of embrittlement (step cooling) was only 5 °C (9 °F)

Fig 25 Influence of microstructure on the temper embrittlement susceptibility of a

chromium-molybdenum-vanadium alloy steel as a function of hardness (a) Actual 50% FATT (b) Change in 50% FATT, F/P, 30% pearlite structure; E, embrittled; NE, nonembrittled Source: Ref 106

ferrite-Embrittlement Predictive Equations. Much of the research on temper embrittlement has concentrated on the influence of composition on susceptibility to temper embrittlement for fixed embrittlement conditions In general, these studies have concentrated on two basic steel grades, nickel-chromium-molybdenum-vanadium and 21

4Cr-1Mo, which are used for rotors and pressure vessels, respectively

In a report on an ASTM study of vacuum carbon deoxidized nickel-chromium-molybdenum-vanadium rotor steels isothermally embrittled at 400 °C (750 °F) for 10,000 h (Ref 107), the shift in FATT (∆FATT) in degrees centigrade was correlated to the impurity content and molybdenum concentration (all in weight percent) by:

∆FATT = 7524P + 7194Sn + 1166As - 52Mo

No significant influence was found for antimony Equation 1 states that phosphorus, tin, and arsenic increased embrittlement, while molybdenum decreased it Also, a phosphorus-tin interaction that decreased embrittlement was observed

A correlation between the 50% FATT and impurity content (J factor) for both nickel-chromium-molybdenum-vanadium

and 21

4Cr-1Mo steels has been demonstrated (Ref 108) The J factor equation is:

where all concentrations are in weight percent

A more detailed correlation has been provided for nickel-chromium steels doped with manganese, phosphorus, and tin (Ref 109) The equation combines the grain-boundary phosphorus and tin concentrations, the prior-austenite grain size, and the hardness level Equation 3 was extended to a nickel-chromium-molybdenum-vanadium steel with hardnesses of

20 and 30 HRC, ASTM grain sizes of No 3 and No 7, and isothermal embrittlement at 480 °C (895 °F) for 6000 h (Ref 110) The resulting equation was:

∆FATT = 4.8P + 24.5Sn + 13.75(7 - GS) + 2(HRC - 20) + 0.33(HRC - 20) (P + Sn) + 0.036(7 - GS)(HRC - 20) (P + Sn)

or service time within a temperature range of 300 to 600 °C (570 to 1110 °F) or because of slow cooling through this range Coarse-grain material is more susceptible than fine-grain material The degree of embrittlement is greater for martensite than bainite and least for ferrite-pearlite However, embrittled martensite is still tougher than nonembrittled bainite, while embrittled bainite is tougher than nonembrittled ferrite-pearlite Additional information may be found in Ref 110, 111, 112, 113, 114

References cited in this section