Hydrogen transport and fracture toughness of case hardened steel

ELSEVIER

Materials

Hydrogen transport

Chemistry

and Physics

38 (1994) 234-242

and fracture toughness of case hardened steel S. C. Lee, W. Y. Wei and L.H. Chiu

Department of Materials Engineering,

Tatung Institute of Technology, 10451, Taipei (Taiwan, ROC)

D. L. Johnson* Department of Mechanical Engineering

Metallurgy Program, University of Nebrada-Lincoln,

Lincoln, NE 68588 (USA)

(Received June 28, 1993; accepted April 2, 1994)

Abstract The purpose of this study is to assess the effects of effective case depth (ECD) and tempering temperature on hydrogen transport and fracture toughness of carburized AKSI 8620 alloy steei. The material was machined into thin discs for permeation and into compact-tension specimens for fracture toughness measurements. The specimens were pack carburized at 930°C and cooled to ambient temperature. The carburized specimens were austenitized at 840°C in a high temperature salt bath, then oil quenched and tempered at various temperatures for one hour. Assessment of hydrogen transport was conducted by the electrochemical permeation technique. Both permeability and effective diffusivity decrease as ECD increases and tempering temperature decreases. Fracture toughness of pre-charged carburized 8620 steel increased with depth of carburization. Fracture toughness of hydrogen pre-charged carburized 8620 steel was minimum at tempering temperatures between about 150-400°C. The results indicated good empirical correlation between fracture toughness and apparent solubility at higher ECD and tempering temperature. While this study shows that correlations exist, there are compositional and microstructural factors to be sorted out before any definitive relationship can be established.

1. Introduction Case hardening by carburizing has long been used to produce high surface strength, better wear resistance and improved fatigue life for critical components. A fairly extensive body of literature exists on the mechanical properties, fatigue and fracture toughness of carburized steels [l-4]. With respect to hydrogen embri~tlement (HE), however, much less info~ation is available on the effect of carburization. Issues [5,6] with regard to hydrogen were discussed at the symposium “Carburizing-Processing and Performance” held July, 1989 in Lakewood, CO, USA. A question arises as to the disposition of the hydrogen during annealing after the hydrogen source is removed. There is concern that hydrogen cannot readily diffuse from the core, thus, remains at lattice or defect sites to promote HE. *Formerly Visiting Professor, Department of Materials Engineering, Tatung Institute of Technology,

1990, (Taiwan, ROC)

0~4-0584~4/$07.~ Q 1994 Elsevier Science S.A. Ail rights reserved SSDZ 0254-0584(94)01373-O

Information now available regarding morphology effects on HE suggest that surface carburized case could either help or hinder resistance to HE. In carbon steels, diffusion and HE susceptibility vary with carbon content, ferrite/carbide morphology and the degree of reversibility of interface traps [7-121. Compressive stresses internally induced in the carburized layer by volume change may have a beneficial effect on HE [13,14]. While good progress has been made in ~derstanding HE behavior in carbon and alloy steeis, extrapolation of these observations may not be valid for more complex carburized microstructures. Thus, the purpose of this work is (1) to assess the effects of effective case depth and tempering temperature on hydrogen transport in carburized AISI 8620 alloy steels, (2) to assess empirically the effects of hydrogen trapping, case depth and tempering temperature on fracture toughness in the same steels. Variables to be examined include carburization depth, tempering temperature and hydrogen charging time.

SC.

Lee et al. I Materials Chemistry and Physics 38 (1994) 234-242

235

2. Experimental Procedures

2.3. Metallographic Analysis

2. 1. Material

Metallographic samples were taken from permeation and CT specimens, ground, polished and then etched with a 4 % Nital solution. The results of the examination were used to confirm case depth obtained by microhardness and to confirm proper response of the samples to heat treatment.

The chemical composition of the AISI 8620 steel samples as analyzed by atomic absorption technique is listed in Table 1. The permeation specimens were cut from round bar and ground to 1.1 mm thick discs. Round steel bars of 80 mm outside diameter were machined into 58 x 60 mm rectangular bars for making the compact-tension (CT) specimens. Specimen geometry is shown in Fig. 1. 2.2. Heat Treating The specimens were case hardened by pack carburizing. The specimens, packed with compounds composed of 70 wt% charcoal and 30 wt% barium carbonate (BaC03) as energizer, were all sealed into a stainless steel crucible with clay. The crucibles were heated to 930°C in an electric furnace for different times and then furnace cooled to room temperature. The carburized specimens were austenitized at 840°C in a high-temperature salt bath, then oil quenched and tempered at 15O”C, 400°C and 700°C for one hour, respectively. The heat-treating process for the pack carburizing is shown in Fig. 2.

TABLE

1. Chemical

Composition

of AISI 8620 Steel (W/O).

C

S

Si

P

Mn

Ni

Cr

MO

0.22

0.016

0.34

0.025

0.52

1.50

0.48

0.17

1.25w

2. 4. Microhardness

Testing

Case depth profiles were determined from Vickers microhardness. The load used for testing was 300 grams. The effective case depth (ECD) was defined as the distance from the surface to a point where the microhardness is 550 HV. 2.5. Electrochemical Permeation Hydrogen transport behavior was studied by the electrochemical permeation technique [7]. The apparatus is shown in Fig. 3. The cathodic side or hydrogen entry cell was galvanostatically polarized at a constant cathodic current density i, = 10 mA/cmZ in 1N HzS04 poisoned with 1 gram/liter of thiourea. The anodic side or hydrogen exit cell was potentiostatically polarized at 300 mV(SCE) in 0.1 N NaOH. Transport parameters were obtained at temperatures between 30 and 45°C. The effective diffusivity was determined from the lag time, tL, at a constant current corresponding to 0.63 x i;. Experimental data were correlated with the permeability, J-L = i; L/nF [mole H m-1s-r],effective diffusion coefficient, Deff= LY6t, [m* s-r],and the apparent solubility, CPP = JML/D,~~ [mole H m-31 where ir is the steady state current density, L is specimen thickness, n is the number of electrons transferred, F is Faraday’s

02

a: Precrack Fig. 1.

Dimensions

length=0.45-0.55

of the CT specimen,

W

B=25 mm, W=48 mm.

CARBONCONTENT(Wt%)

Fig. 2.

Processes

for pack-carburization

TIME

and heat treatment.

236

XC. Lee et al. / Materials Chembtry and Physics 38 (1994) 234-242

2.6. Hydrogen Precharging and Fracture Toughness

0 I N NaOH i ,--,*

I 2 3 4 5

,’

Cathodic portion ofthe cell Anodic portion of the cell Counter electrode Thermometer Gas dispersion tube (h’, gas inlet)

6 Ag/AgCI electrode 7 Working electrode (specimen) 8 Magnetic stirring bar

Hydrogen precharging of the CT specimens was conducted in a 1N H2S04 solution poisoned with one gram/ liter of thiourea at 30°C. Platinum served as the counter electrode. During cathodic charging, the entire CT specimen was immersed into the solution. Specimens were charged with a galvanostat at a constant current density of 4 mA/cm*. The specimen was then removed from the electrolyte and rinsed with distilled water. Within 5 minutes after hydrogen charging was completed, the CT specimens were fracture tested using a Z-Tons MTS dynamic testing machine at ambient temperature. Al1 CTspecimens were fatigue precracked prior to hydrogen charging and final fracturing. The fracture toughness data were calculated following the recommended procedures of ASTM E561-80 [15] and ASTM E399-83 [Xl.

Fig. 3. Schematic diagram of the electrochemical ceil used for permeation measurements.

3. Results and Discussion 3.1. Microstructure and Microhardness Gradients constant and tL is second transient lag time. Normally, initial transients give invalid lag time because initial charging fills irreversible traps first. These traps strongly bind hydrogen at defects such as incoherent interfaces and microvoids. Second transient lag time results from filling of reversibie traps. These traps weakly bind hydrogen at defects such as coherent interfaces, grain boundaries, solute atoms and dislocations and reflect the parameters J-L, Dee and Capp. From these definitions, J-L is determined at steady state whereas Deff is determined from tL. Temperature dependencies are defined by JmL = JZ L exp(-Q~/RT), DeK = D, exp(-Qn/RT) and C,, = C, exp(-AH,,,/RT) where JZ L, D, and C, are permeation, diffusion and solubility pre-exponential constants; Qr, Qn and AHan, are permeation and diffusion activation energy and apparent heat of solution respectively and T is absolute temperature. Combining expressions for JmL and Deff into Carp, it can be shown that AH,, = Qp - Qn The tendency toward embrittlement, as related to transport parameters JmL, De8 and Can,, has been discussed in a general way by Pressourye [13] and applied by Johnson et al. [7, lo] to plain carbon steel. In the present study of carburized 8620 steel, gradient microstructure and trapping behavior is much more complex than that of plain carbon steel. For the purpose of this study, a relatively low Cap,,is thought to minimize the tendency toward embrittlement since transported hydrogen easily drops into widely dispersed irreversible traps rather than crack nucleation sites.

The morphology of the martensite in the case of a carburized specimen can be controlled in several ways. Surface carbon content could be maintained at a level low enough to produce predominantly lath martensite. Another approach would be to reheat a specimen with 1 wt% carbon or more in the case into the austenitecarbide two phase field. Such a treatment would form carbides and reduce the austenite carbon content. The refined microstructure of a specimen reheated into the austenite-carbide field would appear to offer a number of advantages including reduced martensite microcracking, a fine distribution of hard carbide particles, a fine austentitic grain size, and a very fine martensite. Typical microstructures of a reheated carburized 8620 steel tempered at 150°C for 1 hour are shown in Fig. 4. The case reveals high carbon martensite with dispersed spherical carbide and some retained austenite, which is in agreement with Lee [1] and Krauss 121.The core reveals low carbon lath martensite. The grain structure appears to be finer in the case than in the core. Grain size and retained austenite are known to influence toughness properties [l]. They may influence transport properties as well. Hardness profiles of the case-carburized steel as a function of case depth are shown in Fig. 5. The ECD values were found to be 0.13, 0.18, 0.32 and 0.50 mm respectively for steel carburized 15, 30, 60 and 120 minutes. Fig. 5 also shows that steels with shorter carburizing time yield lower surface hardness due to the shallow penetration and a smaller amount of carbon diffusing into the surface.

237

XC. Lee et al. I Materials Chemistry and Physics 38 (1994) 234-242 1000 -

Material: AISI 8620 steel Carburizing time: weee 15 min 30 min A-A-M-A 60 min +HH+0 120 min

0.80

Distance

1. 10

(mm)

Fig. S. Hardness profile gradient as a function of carburi~ation time.

Fig. 4. Microstructure of AISI 8620steelcarburizedat 930”Cquenched from 840°C in oil. and tempered at 150°C. (a) case and (b) core.

.?.2. Hydrogen

Transport

10-4 3.10

1

3.15

1

3.20

1000/T The effect of case depth on hydrogen transport through carburized steels was studied as a function of temperature between 30 to 45°C. The effective diffusivity D,rf and the permeability JmL, obtained at a tempering temperature of 150°C are shown in Figs. 6 and 7, respectively. Thermally activated Arrehenius behavior is indicated in these figures as in the study of Johnson et al. [9]. The data clearly show a decrease in both J-L and Derr as ECD, the overall carbon content of the permeation specimen, increases. The reason for the decrease in J-L and DeB was difficult to assess in the complex tempered martensitic gradient microstructure. However, it appeared that the carburized layer was dominant. As the carbon content is increased, quenching

,

,

3.25

(K-l)

3*30

,I

3.35

Fig. 6. Elffective diffusivity versus inverse temperature as a function of effective case depth, tempered at 150°C.

induces higher densities of lattice imperfections. These imperfections increase the number of reversible hydrogen trapping sites and delay hydrogen diffusion. Thus, Defl decreases with increasing carbon content (ECD) [lo, 111. The effect of tempering temperature on hydrogen transport at ECD = 0.13 mm is shown in Figs. 8 and 9. Reasoning similar to that discussed above may explain the increase in D,@with increasing tempering temperatures above 150°C. Well known structural changes on

238

XC. Lee et al. / Materials Chemistry and Physics 38 (1994) 234-242

50

TemperQMXX) QMIXJ ~AAU

I 3.10

I

AISI 6620,

1

3.15

3.20

ECD= 0 mm ECD=O.i3mm ECD=O.SOmm

I

I

3.25-l

I

3.30

iaI

3.3

1000/T (K )

I

45

I

1

35

t

30

I

25

I

ECD = 0.13 mm As Quenched oooo0 00000 150°C Tempering AAAAA 4OO’C Tempering 00000 700°C Tempering

t

,

3.10

3.15

I

I

1

3*30 1000/T (K-') 3.20

3.25

I

I

i

30

t

3.15

3.20

3.25

3.35

1000/T (K-l)3.30

Fig. 7. Permeability versus inverse temperature as a function of effective case depth, &=lO mA/cmZ, tempered at 1.5OT.

50

i

ECD = 0.13 mm OOQOo As Quenched KIOOOO 150°C Tempering AAAAA 4OO’C Tempering Q0 0 0 0 700°C Tempering

150°C AISI 8620, AISI 8620,

T4,(“C)35

45

I

Fig. 9. Permeability versus inverse temperature as a function of tempering temperature, &=lO mA/cm*, ECD=0.13 mm.

carbide and low-carbon martensite by cementite and ferrite be~nning from 250 to 350°C. According to the study of Craig and Krauss [12], as tempering temperature increased, not only did the width of the martensitic lath and plate increase on the average, but the coarsened structure became more equiaxed. Definite coarsening of the lath martensite structure developed on tempering up to 6OO”C,but above this temperature, the effect was very rapid. In addition to the changes in the matrix of the steels, tempering also significantly changed carbide size, mophorlogy and distribution. The values of J-L and Deff increased as the temper temperature increased due to which reduced the density of reversible traps, structural changes, especially in the specimen tempered at 700°C [lo]. 3.3. Fracture Toughness

3.35

Fig. 8. Effective diffusivity versus inverse temperature as a function of tempering temperature, ECD=0.13 mm.

tempering of carbon steels reduce the number of reversible trapping sites and JmL and I),@ increase. From Speich [17], three distinct stages were identified: Stage I: The formation of a transition carbide, epsilon carbide, and lowering of the carbon content of the matrix martensite (100 to 250°C). Stage II: The transformation of retained austenite to ferrite and cementite (200 to 300°C). Stage III: The replacement of the transition

Fracture toughness as a~nction hydrogen pre~harging time and case depth was determined by using compacttension specimens. The Pmw/Po value must be lower than 1.1 in order to justify the validity of plain-strain KIc values [15]. Based on this criteria, only plane-stress K, values were valid and are shown in Tabies 2A, 2B, 2C. Fig. 10 shows fracture toughness versus hydrogen precharging time for carburized specimens with ECD = 0.32 mm, tempered at 150°C. Precharging up to one hour had no effect on K,. However, K, decreased for precharging time longer than one hour, presumably due to hydrogen trapping. This suggests that the carburized layer may, in some way, delay the onset of hydrogen

239

S.C. Lee et al. I Materials Chemistry and Physics 38 (1994) 234-242 300

TABLE 2A. Kc as a Function of Hydrogen Precharging Time, ECD = 0.32 mm, Tempered at 150°C. Precharging time

pmxpQ

c

ECD = 0.32 mm Temper150°C

Kc (MPa ml’*)

(hr) 0 0.33 1 2 2 4 4

1.428 1.979 1.360 1.315 1.302 1.283 1.273

233.5 233.3 235.0 180.1 183.7 171.2 166.0

TABLE 2B. Kc as a Function of ECD, Tempered at 15O”C,2 Hours Hydrogen Precharging. ECD

Ptna~~~Q

Kc (MPa m1j2)

(mm) 0 0 0.18 0.18 0.32 0.32 0.40 0.40

1.214 1.274 1.408 1.410 1.315 1.302 1.410 1.380

131.9 146.6 191.0 165.9 180.1 1x3.7 192.9 168.6

TABLE 2C. KC as a Function of Tempering Temperature, ECD=0.32 mm, 2 Hours Hydrogen Precharging. Tempering Temperature (“C)

prnax/pQ

Fig. 10. Fracture toughness versus hydrogen ECD=0.32 mm, tempered at 150°C.

precharging

time.

TABLE 3A. Cant, as a Function of ECD*

(mm)

J,L x lo9 (mole H m-rs-l)

D & x 10’0 ( m2s-t )

c aPP (mole H m”)

0.00 0.13 0.50

8.3 6.4 1.3

2.00 1.00 0.23

41.5 64.0 56.3

ECD

*Tempering Temperature-150°C Transport-35°C

Kc (MPa rnln) TABLE 3B. Cap,, as a Function of Tempering Temperature *

As-quenched As-quenched 1.50 1.50 400 400 700 700

1.317 1.414 1.315 1.302 1.294 1.283 1.675 1.3%

288.4 343.8 180.1 183.7 183.6 172.0 335.1 340.9

embrittlement. The relationship between ECD and the variables D,s, J-L, C,, and K, are shown in Figs. lla, b, c and d. The relationship between tempering temperature and the same variables are shown in Figs 12a, b, c and d. An initial increase in both C,, and K, with increase in ECD appears to be inconsistent with earlier discussion since more hydrogen would be available to drop into crack nucleation sites. Fig. lld shows a trend which is similar to that reported by Lee and Ho [l] for non-precharged specimens. K, values clearly increase as a function of effective case depth. Since the case structure is primarily

Tempering Temp.(“C) None 150 400 700

J_L x lo8 (mole H m-*s-r) 1.17 0.64 1.25 1.45

D,@ x lOI* (m2s-‘)

c =PP (mole H m”)

1.20 1.00 2.00 6.00

97.5 64.0 62.5 24.2

*ECD-0.13 mm Transport-3Y’C

high carbon tempered martensite, it can be postulated that the static fracture toughness of the high carbon tempered martensitic structure in the case is tougher than that of the low carbon tempered martensite in the core. Thus, increasing the volume fraction of the case improved the fracture toughness of the surface-hardened steeis. The grain structure of the case was finer than that of the core. Since fine grain structure is known to have better toughness than coarse grain, the greater fine grain areas of the deeper case might have contrib-

240

XC. Lee et al. I Materials

Chemistry and Physics 38 (1994) 234-242 80.0

Diffusivity-35°C

t

Solubility-35°C Temper- 150°C

0

&WV? V I

I

0.2

0.3

I

0.0

0.1

(4

ECD

I

(m:)

I

0.5

30.0

0.6

1

01.0

I

I

CI

I

0.3

I

0.4

ECD (mm)

I

0.5

0.6

250.0

Permeability-35°C Temper- 150°C

Temper150°C Hz Precharge-2

hours

i

I

0.0

I

0.1

0.2

I

ECI?‘( rn:)

I

I

0.5

I

0.6

I

rz

50.00 k

(b) Fig. 11.

0.2

(4

I

h

I

0.1

(a) Effective Table 2B)

diffusivity

I

o 1

I

0.2EC;.3cm;I

1

I

0.5

0.6

(d) vs ECD, (b) Permeability

vs ECD, (c) Apparent

uted to the improved toughness as well. The finer grain structures can be clearly observed metallographitally in Fig. 4(a). However, at higher ECD, Carp decreases. This decrease is consistent with increase in K, assuming that additional irreversible trap sites become available when ECD increases. It can be concluded that the extent of case depth and higher-carbon tempered martensitic structure are the dominating factors in improving the K, fracture toughness of the carburized steels. Hydrogen precharging does not appear to effect this dominance at higher ECD. An initial decrease in both C,, and K, with increasing tempering temperature again appears inconsistent

solubility

vs ECD, (d) & vs ECD. (Data shown in Table 3A and

with earlier discussion since less hydrogen would be available to drop into crack nucleation sites. This may be due to the difference in fracture behavior between as-quenched and tempered structure involved the hydrogen reaction. A suitable explanation will have to await a more systematic study. However, at tempering temperatures above -15O”C, K undergoes a reversal consistent with Carp. In contrast to the results of Lee and Ho [l], the specimen tempered at 700°C had the highest K value in the precharged state. According to Graig and Krauss [12], as the specimens are tempered at successively higher temperatures, the fracture mode in hydrogen changes from intergranular to

S.C. Lee et al. I Materials

241

Chemistry and Physics 38 (1994) 234-242

Solubility-35°C

I

0

1

200

Tempering (a>

I

400

20.0

I

600

Temperature

,

0

I

I

200

Tempering (b)

I

400

400

E

I

600

Temperature

E (‘t

Oo

I

200

TempT;;ng

Fig. 12. (a) Effective diffusivity vs tempering temperature, (b) Permeability vs tempering temperature, temperature, (d) Kc vs tempering temperature. (Data shown in Table 3B and Table 2C.)

transgranular. The specimen tempered at 700°C resulted in a high P,,,, /Po value (1.675) i.e., ductile fracture mode, can be explained the importance of low apparent hydrogen solubility. It is clear from this discussion that while empirical correlation exists between hydrogen transport and fracture toughness, there are many compositional and microstructural factors to be sorted out before any definitive relationship can be established.

1

600

Temperature

ECD-0.32 mm H2 Precharge-2

Permeability-35°C ECD-0.13 mm

'0

200

Tempering (4

(O?

,

I

400

kl t

(“C

1 hours

I

600

Temperature

o 800

( C)

(c) Apparent solubility vs tempering

4. Conclusion It is concluded from this study that both permeability and effective diffusivity of hydrogen decreases in carburized steels as ECD increases and tempering temperature decreases. Fracture toughness of precharged carburized 8620 steel increased somewhat with depth of carburization. Fracture toughness of hydrogen pre-charged carburized 8620 steel was minimum at tempering temperatures between about 150-400°C. An increase in fracture toughness at higher case depths and tempering temperatures is consistent with a corresponding decrease in apparent solubility.

242

S.C. Lee et al. / Materials Chembtry and Physics 38 (1994) 234-242

Acknowledgments The authors are grateful for the financial support of this research by the National Science Council, Republic of China, under contract No. NSC 79-040%E036-12.

References 1. SC. Lee and W.Y. Ho, Metall. Trans. A. 20A (1989) 519. G. Krauss, Metall. Trans. A, PA (1978) 1527. D.V. Doane, .f Heat Treuring, 8 (1990) 33. D.E. Diesburg and G.T. Eldis, Metall. Trans. A, PA (1978) 1561. G. Krauss, Carb~ri~~ng-Process~g and Performance, ASM Lakewood, July 1989. 6. H. Streng, C. Razim and J. Grosch, in G. Krauss (ed), Carburizing Processing and Performance, ASM Lakewood, 1989, p. 311. 2. 3. 4. 5.

7. H.W. Jeng, L.H. Chiu, D.L. Johnson, and J.K. Wu, Metall. Trans. A, 2IA (1990) 3257. 8. S.X. Xie and J.P. Hirth, Corrus~on-BAKE, 38 (1982) 486. 9. D.L. Johnson and J.K. Wu, J. Mat. for Ener. Sys., 40 (1987) 402. 10. D.L. Johnson, G.Krauss, J.K. Wu, and K.P. Tang, Metull. Trans. A, 18A (1987) 717.

11. S.L.I. Chan, H.L. Lee and J.R. Yang, in N.R. Moody and A.W. Thompson (eds.), Prof. 4th Int. Conf on Hydrogen Effects on Mater~ul Behavior, TMS, Warrendale PA, 1990, p. 45. 12. B.D. Craig and G. Krauss, Metall. Trans. A, 1lA (1979) 1799. 13. G.M. Pressouyre, Metall. Trans. A, 14A (1983) 2189. 14. G.M. Pressouyre, Metall. Trans. A, 1OA (1979) 1571. 15. ASTM E561-80, Annual Book ojASTM Standards, Part 10, ASTM, Philadelphia, PA, 1980, p. 656. 16. ASTM E399-83, Annual Book ojASTM Standards, Part 10, ASTM, PhiIdelphia, PA, 1984, p. 518. 17. G.R. Speich, Metals Handbook, Vol. 88th ed., ASM, Metals Park, Oh., 1973, p 202.