Materials' Science and Engineering, A 142 ( 1991 ) 235-243 235

Effects of hydrogen pre-charging for quenched-and-tempered AISI 1520 steels containing boron

A. Chatterjee* AdTech Systems Research Inc'., Dayton, OH 45432 (U.S.A.)

R. G. Hoagland and J. P. Hirth Department of Mechanical and Materials Engineering, Washington State University, Pullman, WA 99164 (U.S.A.)

(Received December 28, 1990; in revised form March 6, 1991 )

Abstract

A quenched-and-tempered AIS1 1520 steel containing boron, as well as similar steels with no boron additions, were tested in plane strain conditions after hydrogen pre-charging at a range of hydrogen fugacities. Above a critical charging density that varied with heat treatment and testing parameters, the strain at fracture for the boron-containing steels dropped to zero and the fracture surface became mixed intergranular and quasi-cleavage, with the fraction of the former increasing with increasing charging current density. No such transition was observed for the steels without boron. A model for the dynamic drag of a hydrogen atmosphere by a propagating crack is proposed to account for the observations.

1. Introduction

Although hydrogen embrittlement has been studied in a large number of metals and alloys, there has been no single theory to explain all the observed effects of hydrogen-induced degrada- tion in steels. In the presence of hydrogen, high strength steels are very sensitive to the presence of foreign particles and impurities, and the frac- ture process is often brittle with cleavage and quasi-cleavage fracture surfaces. When large amounts of segregated impurities are present at the grain boundaries, hydrogen has also been observed to promote the decohesion of grains and cause intergranular fracture. For such frac- ture the presence of grain boundaries, which generally strengthen a metal, is an inherent source of weakness. Briant and Banerji [1] have indicated that among the circumstances inducing inter- granular fracture the important ones are the pres- ence of grain boundary phases, segregation of impurities at grain boundaries and the action of various environments.

*Present address: The Materials Development Labora- tory, Allison Gas Turbine Division, General Motors Cor- poration, Indianapolis, IN 46260, U.S.A.

Hydrogen has a strong tendency to segregate at structural defects such as edge dislocations and grain boundaries. This has been shown indirectly in a number of permeation experiments [2] and directly in the tritium autoradiography experi- ments of Aucouterier [3]. Also tritium auto- radiography of the same heat of the AISI 1520 steel as studied here, in the spherodized condi- tion [4], indicated that carbides are weak traps for hydrogen. Hydrogen segregation controls the properties of hydrogenated materials in many cases, common examples being hydrogen diffu- sion, hydrogen-induced cracking and most importantly the mechanical behavior in many engineering materials exposed to hydrogen environments. The strong tendency of hydrogen to interact with dislocations and grain boundaries influences the solid solubility and the mobility of hydrogen and consequently modifies the embrit- tlement process. Also intergranular segregations of alloying elements and other impurities modify the susceptibility of the material because of the interactions of the impurity and the hydrogen atoms. The existence of fracture at prior austenite grain boundaries in quenched-and-tempered alloy steels, which correlates with cosegregation effects of hydrogen and group V and VI impuri-

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ties, is the clearest evidence of the influence of hydrogen on the grain boundaries [5]. Segrega- tion models to explain this phenomenon have been proposed by Rice and coworkers, reviewed in ref. 6.

For boron-containing steels, Druce [7], Mait- pierre et al. [8] and Inoue et al. [9] have all shown the detrimental effects of borocarbide segrega- tion in quenched-and-tempered steels after hydrogen charging. Fine precipitate particles closely spaced along the grain boundaries can block the movement of dislocations and thereby increase the dislocation density and the pile-up stresses near the crack tip. This local hardening process not only decreases plastic relaxation at the crack tip but also promotes the local accumu- lation of hydrogen (adsorbed in the strain fields and in the cores of dislocations), which is in turn important in enhancing decohesion and conse- quently fracture.

In this work, the interaction of hydrogen and boron, present in the elemental form or as boro- carbides, to cause partly intergranular fracture was examined. A simple model is proposed to account for stress-assisted diffusion of hydrogen to the crack tip together with the drag of hydro- gen by the moving dislocations during the defor- mation process. On the basis of this model, a phenomenological explanation of the influence of charging current densities and strain rates in causing partly intergranular fracture in the quenched-and-tempered AISI 1520 steel is sug- gested.

2. Experimental procedure

The primary steel investigated in this research was a laboratory heat of a boron-containing AISI 1520 steel. The heat was made at the U.S.

TABLE 1

Chemical compositions of the steels studied

Element Amount (wt.%) in the following steels

AISI 1520B AISI 1522 AISI 1020

C 0.18 0.24 0.18 Mn 1.26 1.23 0.61 P 0.003 0.012 0.009 S 0.003 0.008 0.036 Si 0.3 0.25 0.25 AI 0.041 0.049 -- B 0.0008 0.0002 0.0001 Fe Balance Balance Balance

Steel Research Laboratory and hot rolled to plate 13 mm thick. The chemical composition of the steel is given in Table 1. Two other steels (AISI 1522 and AISI 1020) were also studied to some extent to compare the effect of boron and other impurities on hydrogen-induced degradation. These steels were supplied by the LTV Corporation in the form of 40 mm x 40 mm square hot-rolled bars (AISI 1020) which were sliced to 13 mm plate, and 6.5 mm hot-rolled sheet (AISI 1522). The AISI 1020 and 1522 steels contained only trace amounts of boron. The compositions of the various steels are given in Table 1.

The steels were austenitized at 1000 °C for 1 h and immediately quenched in an iced 5% alky- lene glycol-water solution so as to obtain a fully martensitic structure. Subsequently, quenched plates were tempered at different temperatures. The prior austenite grain size was about 100 #m. Clausing-Hill plane strain samples [10] were machined from these heat-treated plates. The samples were then mechanically polished to 0.25 #m diamond paste. The direction of final polish- ing was parallel to the rolling direction which was the same as the direction of the applied tensile stress. Mechanical polishing was chosen over electrolytic polishing because of the adverse effects of pit formation accompanying the use of electrolytic techniques [11].

The charging set-up was similar to that in ref. 12. Specimens masked with acid-resistant paint except for the gauge section were pre-charged in a solution of i N H2SO a with 1 g thiourea 1- l as a hydrogen recombination poison. The solution was deaerated by bubbling with high purity deoxygenated nitrogen for at least 8 h prior to testing. The charging current density was varied between 6 and 200 A m-2. This range of charg- ing current density corresponds roughly [13] to a fugacity of 200 MPa to 2 GPa and a hydrogen concentration range from 0.2 to 0.4 tool m -3 respectively. All specimens were charged for 2 h except for those where the effect of hydrogen was studied as a function of the charging time (see Tables 5 and 6 later). Hydrogen-charging experi- ments were done at room temperature. At the end of charging, the samples were quickly removed from the charging cell, rinsed in flowing water to remove the residual acid drops and dried in acetone.

Subsequently the samples were subjected to mechanical testing in an Instron testing machine

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within 2 min of charging. Most of the mechanical tests were done on a screw-driven Instron model TT. A few of the tests were also done on a servo- hydraulic Instron model 1322. The cross-head speeds as well as the actuator movement were 8.4 #m s- l which corresponded to a strain rate of 1.33 x 10 -3 s-~. All mechanical test results pre- sented are the average of at least three tests. The variation in the results was _+ 10%.

The fracture surfaces were examined using a scanning electron microscope. Since some of the quenched-and-tempered specimens fractured intergranularly, at the prior austenite grain boundaries, scanning Auger microscopy (SAM) was conducted to detect the segregated impuri- ties.

3. Results

The major results are presented in Tables 2-4. The low-boron-containing steels failed in a ductile dimple and quasi-cleavage manner under all conditions tested. There was some degrada- tion induced by hydrogen charging, mainly re- flected in a decrease in fracture strains. These results parallel those for the AISI 1020 steel in the spheroidized condition [14] where hydrogen produced moderate degradation in tensile behav- ior even under dynamic charging conditions,

unlike results on higher carbon steels where the degradation was greater [12, 15]. In the present low-boron-containing steels, the fracture mode was mainly ductile dimple rupture in all cases, with some increase in quasi-cleavage areas after hydrogen charging. The latter finding is consis- tent with the effect of hydrogen in promoting shear instability [12, 16].

In contrast, the AISI 1520 steel showed a complete brittle fracture behavior for charging current densities greater than or equal to 20 A m -2. Also, the fracture appearance was of a mixed quasi-cleavage-intergranular nature with the area fraction of intergranular fracture increas- ing with increasing charging current density (Table 2 and Fig. 1 ). As discussed in detail earlier [17], SAM studies of intergranular surfaces of a sample of AISI 1520 fractured in situ showed boron segregation to prior austenite grain bound- aries in an amount of 2 at.%. The boron was either elemental or in the form of fine boro- carbide clusters.

Some auxiliary tests were also performed on the AISI 1520 steel as indicated in Tables 5-8. The degradation increased with increased charg- ing time as shown in Table 5. However, as shown in Table 6, the degradation is completely revers- ible under the charging conditions studied. In the permeation experiments of Xie and Hirth [13], it

TABLE 2

Mechanical properties for the quenched-and-tempered (400 °C) AISI 1520 steel as a function of the charging current density for a charging time of 2 h

Current density Yield strength Ultimate tensile Uniform strain Fracture strain Fracture (Am 2) (MPa) strength (MPa) e, e I appearance

0 1040 1250 0.05 0.20 DD 6 990 1140 0.049 0.049 DD + QC

20 ~' 431 0 0 QC + 15% IGF 50 ~' 425 0 0 QC + 50% IGF

170 " 42(1 (I 0 Q C + 85% IGF

DD, ductile dimple rupture; QC, quasi-cleavage; IGF, intergranular fracture. ~'Fractured prior to yielding.

TABLE 3

Effect of hydrogen on the mechanical properties of quenched-and-tempered AISI 1020 steels as a function of the charging current density for a charging duration of 2 h

Current density Yield strength Ultimate tensile Uniform strain Fracture strain Fracture (A m 2) (MPa) strength (MPa) e U e I appearance

0 831 1130 0.06 0.13 DD 20 825 870 0.04 0.04 DD + QC

DD, ductile dimple rupture; QC, quasi-cleavage.

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TABLE 4

Effect of hydrogen charging on the mechanical properties of quenched-and-tempered A ISI 1522 steel as a function of the charging current density for a charging time of 2 h

Current density Yield strength Ultimate tensile Uniform strain Fracture strain Fracture (A m 2) (MPa) strength(MPa) e u e~ appearance

() 725 826 0.07 0.19 DD 20 740 841 0.07 0.15 DD + QC 80 723 856 0.05 0.12 DD + QC

DD, ductile dimple rupture; QC, quasi-cleavage.

Fig. 1. Montages of the fracture surfaces of quenched-and-tempered (400 °C) AISI 1520 steels after hydrogen charging for 2 h at a current density of (a) 20 A m 2, (b) 80 A m- 2 and (c) 170 A m :.

has been demonstrated that the permeation current reaches steady state values in minutes. The time needed for the ratio of the transient flux to the steady state flux of hydrogen to be 0.5 is roughly 10 min for a specimen 2 mm thick (simi- lar to our Clausing-Hill specimen) and charged at a current density of 50 A m 2 [13]. Also, as shown in Table 7, pre-straining in tension after tempering and prior to hydrogen charging increased the fracture stress, with the increment of increase diminishing with increasing pre-strain while the fracture mode remained brittle, ef = 0. A few experiments were carried out for other tempering temperatures. The results (Table 8)

suggest that the degree of embrittlement de- creases with decreasing strength levels. However, the tempering temperature of 260 °C is in the range of tempered martensite embrittlement, which may also influence the hydrogen fugacity needed for complete embrittlement.

Finally a notched Charpy sample of AISI 1520 was charged at 200 A m -2 for 2 h and fractured by impact. A montage of the fracture surface for this high strain rate test (Fig. 2) shows no intergranular fracture. The same steel when charged at 175 A m -2 for 2 h revealed over 75% intergranular fracture when deformed at a strain rate of 1.3 × 10 -s s -l (Fig. l(c)).

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TABLE 5

Effect of charging duration on the mechanical properties of quenched-and-tempered (400 °C) AISI 1522 steel as a function of the charging current density for a charging current density of 6 A m 2

Time Yield strength Ultimate tensile Uniform strain Fracture strain Fracture (h) (MPa) strength (MPa) e u ej appearance

2 99(1 1140 0.048 0.09 DD + QC 10 949 1050 0.02 0.1)2 DD + QC 24 758 89(1 0.01 0.01 DD + QC

TABLE 6

Effect of hydrogen outgassing for 1 month after charging of quenched-and-tempered (400 °C) AISI 1520 steel

Current Time Yield strength Ultimate tensile Uniform strain Fracture strain Fracture density (h) (MPa) strength (MPa) eo ej appearance (A m 2)

6 2 1100 1190 0.045 0.195 DD 6 24 1109 1250 0.05 0.2 DD

2(I 2 1040 1190 0.04 0.2 DD

DD, ductile dimple rupture.

TABLE 7

Effect of pre-straining and subsequent hydrogen charging at 20 A m 2 for 2 h on a quenched-and-tempered (400 °C) AISI 1520 steel

Pre-strain Fracture stress Fracture strain (%) (MPa) ef

0 435 0 1.5 1050 0 2.5 935 0 5.11 881 0 6.5 521 0

4. Dynamic model for hydrogen diffusion

The effects of charging current density and strain rate imply that the hydrogen is effective through a dynamic effect on crack propagation. A dynamic model for hydrogen-induced intergran- ular cracking has been proposed by Kameda and Jokl [18]. While it has the correct features to explain the results, it is qualitative. Hippsley et al.

[19] have treated the solute diffusion to a crack tip, relevant to the present problem, but in the slow-velocity drift-controlled case only. An exact solution for a moving potential well with a solute atmosphere is quite complex [20, 21] involving Bessel and Mathieu functions, and the crack problem would be even more complicated.

Hence, to examine the consequences of a dynamic treatment but to avoid the difficult mathematical complexities of a full two-dimen- sional treatment, we simplify the problem to that of a moving one-dimensional potential well [22]. Once the overall features of the problem have been described, the two-dimensional result of Hippsley et al. [19] can be used to estimate the amount of hydrogen in the dynamic atmosphere.

The static interaction energy between a hydro- gen atom and an elastic mode ! crack is [23]

I + V K I Avcos (1) W = 3 r 1/2

where v is Poisson's ratio and is taken to be 0.3, K t is the stress intensity factor, r is the distance from the crack tip, v ( = 3 x 10 -30 m 3 atom -I ) is the molar volume expansion produced by a hydrogen atom and the average value of cos(0/2) is 0.5. In the uncharged case, K I = 50 MPa m 1/2 has been assumed to represent the applied stress intensity for these steels.

The solution to the diffusion equation for the moving one-dimensional potential well is of the form [221

C = K e x p k T (2)

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TABLE 8

Mechanical properties of AISI 1520 steel quenched and tempered at 260 and 600 °C for 2 h

Current Tempering Yield Ultimate Uniform Fracture Fracture density temperature strength tensile strain strain appearance (A m - 2) (°C) (MPa) strength e, ef

(MPa)

() 260 l 140 1420 0.039 0.12 DD + QC 6 260 ~' 1040 0 0 IG + QC

20 260 a 320 0 0 IGF (I 600 515 714 0.118 0.26 DD

20 600 530 731 0.124 0.24 DD

DD, ductile dimple fracture; QC, quasi-cleavage; IGF, intergranular fracture. "Fractured prior to yielding.

Fig. 2. Montage of the fracture surfaces of quenched-and-tempered (400 °C) AISI 1520 steel charged at ! 75 A m-2 for 2 h and fractured under impact loading.

where W has been defined earlier, vc is the crack or well velocity, D is the diffusivity o f hyd rogen in the lattice and x is the distance f rom the well. Since the equat ion involves an exponential te rm involving W/kT and Gx/D, the effects of bo th the terms are equal when they are of equal magni- tude. Drift to the crack tip and consequen t embr i t t lement are caused by hyd rogen accumu- lating within a distance given by the relat ion

w=vcx* (3) kT D

Substituting for W f rom eqn. ( 1 ) and setting r = x, we have

3 K I A v - - - - vex* x' /2kT ~x*/ D (4)

where x* indicates the "size" of the crack field. Solving for x* after substituting the values of various parameters along the est imated crack velocities of 0.15 m m s - ~ and D = 2 × 10 - 9 m 2 s - t we find that x* = 1.78 × 10 -s m.

Hence , when the crack is moving at velocities o f 0.15 m m s -~, the hyd rogen a toms which are 17.8/~m away f rom the crack tip should reach the tip and contr ibute to decohesion. At higher crack velocities this dis tance becomes lower and so there may be a critical velocity and distance which in turn relate to the min imum amoun t of hydrogen that can arrive at the crack tip and still cause decohesion. A lower b o u n d for this critical velocity is the b reakaway velocity, analogous to that for dislocations [22]. As shown for the dis- locat ion case, the condi t ion for b reakaway is roughly when x* = fl = D/vo, i.e. when the region

over which diffusion to the crack tip shrinks to the size of the elastic crack tip field. Guided by the dislocation result [21], a better estimate is given by fl = 4D/v c to account for the actual two- dimensional geometry. Thus in this case the criti- cal crack velocity is given by G=4D/ f l=0 .2

-1 mm s Hence i't is possible for AISI 1520 steel that

slow crack growth at velocities less than about 0.2 mm s -~ can attract sufficient amounts of hydrogen to the moving crack tip for intergranu- lar fracture, while faster crack velocities would produce a dimple rupture fracture. This is pre- cisely the behavior observed in the present work. Moreover the load-displacement record together with a visual estimate indicated that the crack velocity when the samples failed intergranularly was about 0.15 mm s- ~ consistent with v < G and a drift-controlled diffusion of hydrogen to the moving crack tip in the present study.

In the drift regime, the result of Hippsley et al. [19] can be used to estimate the magnitude of the excess hydrogen atmosphere in the elastic field of the crack:

{(2) N(t) =2"5 9~

l/2 (l+v)Kt DJ 4Is Co~-}[32r11/5

(5)

When t>x*2/D, the crack should approach steady state velocity and t in eqn. (5) should be replaced by x*2/D. Substituting with c 0 = 0.3 mol m-3 (corresponding to a charging current density of 20 A m -2) and other values for the various variables we obtain N(t)= 4.06 x 1014 atoms m- 1 This is of the order of magnitude of the static atmosphere for a crack [23] and indicates that there can be near saturation at the crack tip. In theory this suggests that there is sufficient hydrogen at the crack tip to provide continued embrittlement during crack propagation.

5. Discussion

The above model shows that appreciable hydrogen can be dragged along with a moving crack so that intergranular fracture can propagate with a continued role of hydrogen-induced degradation. Equation (5) specifically shows that the degree of degradation should increase with increasing hydrogen fugacity, consistent with the transition to intergranular fracture and the increasing fraction of intergranular fracture with

241

increasing charging current density. Also eqn. (5) suggests that an increasing strength level (through KI) should enhance embrittlement, consistent with Tables 2 and 8. The velocity is coupled to the intergranular area fraction under a given set of conditions. If the intergranular fraction is too large, the crack velocity will decrease and the crack will blunt, becoming a part of the ductile fraction. The reverse would occur if the inter- granular fraction were too small. Hence, under a given set of governing factors, N c and the inter- granular fraction will be fixed.

However, the mere presence of hydrogen is insufficient to reduce the cohesive strength of the boundary and thereby cause intergranular frac- ture as indicated by the absence of an intergranu- lar transition in AISI 1020 and 1522 steels; boron in the AISI 1520 steel must have some effect. Complex borocarbides of the form M23(B, C)6 have been found by other investigators [7-9] to be primarily responsible for hydrogen-aug- mented intergranular fracture even in uncharged steels although for one case in the uncharged condition the intergranular mode was ductile with the dimples forming at the borocarbide particles located at the grain boundaries [9]. In the present AISI 1520 steel, however, there was no inter- granular rupture in the uncharged specimens. Also there was a critical current density of 16 + 4 A m -2 below which there was no intergranular fracture.

Hence there appears to be a cooperative effect between boron and hydrogen in producing embrittlement. One possibility is a direct effect of boron and hydrogen adsorbed at the prior aus- tenite grain boundaries to enhance grain bound- ary decohesion, analogous to the combined effect of hydrogen and group V and VI elements for temper-embrittled steels [24]. Secondly, if fine borocarbides are present, the hydrogen would interact to promote decohesion at the interfaces, contributing to a decrease in the average energy needed to separate the boundary portions at such particles with respect to the energy in the absence of particles.

The effect of pre-straining is second order and dual in nature. One consequence is a break-up of the prior austenite grain boundary by local slip, a factor that produces an initial large increase in the fracture stress, while still resulting in brittle frac- ture with ef = 0 (Table 7). The effect diminishes with increasing pre-strain. This trend is asso- ciated with an increase in internal stress with

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increasing pre-strain, known eventually to cause damage in the presence of hydrogen [13].

The degradation produced by increased charg- ing time in Table 5, and for charging current densities below 20 A cm -2 in Table 2, is un- related to intergranular fracture. Similar results have been found for round-bar tension tests after hydrogen charging. Even though the hydrogen effects are reversible in the sense that the prop- erties are recovered after outgassing, increased charging time apparently produces reversible damage in the material. As suggested by Park and Thompson [16] the damage may be in the form of interfaces that decohere in a time-dependent manner.

Therefore the overall view that emerges is that the propensity of the steel to fail by the inter- granular fracture mode seems to be linked with (a) the hydrogen concentration, (b) the amounts and nature of the segregated impurities, (c) the crack velocities, (d) the diffusivities of hydrogen and (e) the strength and microstructure of the material. The interdependence and interactions of each of the five factors seems to be an impor- tant cause for hydrogen-induced intergranular fracture. Also the effect of segregation would modify the implication of the work of Thompson and Bernstein [25] that showed the tempered martensitic structure offering the best resistance to hydrogen embrittlement, the proviso being that no damaging segregant or precipitate be present at the prior austenite grain boundaries.

6. Summary and conclusion

Plane strain tension specimens of AISI 1520 steel were studied under a variety of heat treat- ments and charging conditions. In the specimens quenched and tempered at 260 and 400 °C, the fracture mode changed from ductile to intergranular after hydrogen charging above a critical current density. The amount of inter- granular fracture increased as the charging current density was increased. Auger analysis after in situ fracture revealed significant amounts of boron at the prior austenite grain boundaries. The. grain boundaries themselves were not inherently weakened by either the segregated boron or hydrogen to cause intergranular frac- ture. In AISI 1020 and 1522 steels, with similar compositions, but no boron addition, no inter- granular fracture could be seen after heat treat- ment and exposure to t h e same hydrogen

environments. Evidently, intergranular fracture was caused by the interactions of segregated boron, the concentration of the available hydro- gen as well as the testing parameters, crack propagation velocity and the strength and struc- tural characteristics of the material. A model has been proposed for the stress-assisted diffusion of hydrogen to the crack tip which is consistent with the role of crack propagation velocities in the process of intergranular fracture and also explains the observed transition from intergranu- lar to transgranular fracture in quenched-and- tempered AISI 1520 steel. The large sensitivity of the boron-containing steel to degradation by hydrogen has implications for design.

Acknowledgments

This work was a part of the doctoral disserta- tion of one of the authors (A.C.) at The Ohio State University. The authors are grateful for the support of this research from the National Sci- ence Foundation under Grants DMR 8361120 and 8813972. They are also grateful to I. M. Bernstein and the U.S. Steel Research Labora- tory for arranging to supply the AISI 1520 steel, to LTV Corporation for the AISI 1522 steel and to I. M. Bernstein and A. W. Thompson for many helpful discussions.

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