Hydrogen effects on deformation — Relation between dislocation behavior and the macroscopic stress-strain behavior

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ScriptaMetallurgicaet Materialia,Vol. 31, No. 2, pp. 149-153, 1994 Copyright©1994ElsevierScienceLtd Printed in the USA. All rights reserved 0956-716X/94 $6.00 + 00

HYDROGEN EFFECTS ON DEFORMATION - RELATION BETWEEN DISLOCATION BEHAVIOR AND THE MACROSCOPIC STRESS-STRAIN BEHAVIOR H. K. Birnbaum Materials Research Laboratory and Department of Materials Science and Engineering University of Illinois at Urbana-Champaign Urbana, IL 6180

(Received March I, 1994) (Revised March 22, 1994) Introduction Direct observations of the effects of H on dislocation mobility, carried out using an environmental cell TEM technique (1), have shown that the presence of H in solid solution increases the dislocation velocities at constant stress for edge, screw, and mixed dislocations and for isolated dislocations and dislocations in "tangles". This increased dislocation velocity due to H has been observed in fcc (2), bcc (3), and hcp (4) systems and in relatively "pure" materials, in solid solution alloys, in precipitation strengthened alloys (5), and in g a m m a prime strengthened alloys (6, 7). "Hydrogen softening" or increased dislocation mobility due to hydrogen has been shown to result from "elastic shielding" which minimizes the elastic interaction between dislocations and obstacles (8, 9). In contrast to the microscopic observations, many stress-strain measurements show an increase in the measured stress for plastic deformation when H is used as an alloy element (10, 11). However, Kimura et al (12-14)have shown that careful electrolytic charging of high purity Fe with H gives rise to very significant decreases in the flow stress over a wide range of temperature. Deformation with simultaneous electrochemical charging of H has been reported to cause both softening and hardening of steels (10). The method is often complicated by introduction of dislocations during the charging which masks the intrinsic H effects. This complexity was avoided in the studies by Kimura et al (12-14). Hydrogen softening was also observed during plasma charging of Fe (15) and in high supersaturation solid solutions of H in A1 (16). Stress relaxation and other techniques which examine thermally activated dislocation mobility have shown that the rate of stress relaxation increases in the presence of H (17). These results have been interpreted in terms of H decreasing the activation barriers which impede dislocation motion. In making comparisons between differing experiments, several factors need to be considered. Hydrogen has a high mobility in many of the systems studied. As a consequence, H atmospheres can move with the dislocations and affect their interactions with obstacles, if the dislocation velocity is less than the atmosphere breakaway velocity (18). Furthermore, the value of B/kT, where B is the binding enthalpy to the dislocations, must be sufficiently large to form a H atmosphere at the dislocation - in the absence of a mobile H atmosphere the H solutes will behave as any other weakly interacting solute. As a consequence of these factors, hydrogen will affect dislocation motion in most systems at temperatures somewhat below room temperature (where the atmospheres can form), and at low strain rates (where the atmospheres are mobile).

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Most experiments which have shown d e c r e a s e s in the flow stress and increased dislocation mobility due to H meet the above criteria. In some, but not all, cases which report increases in the flow stress due to H, the strain rates have been relatively high. Another factor which must be considered in comparing experiments, is whether the alloy forms hydrides under the conditions of the experiment. Formation of hydrides (as in the systems based on Ni, Ti, stainless steel, group Vb elements, and Mg) would generally lead to hardening due to the very large AV/V which accompanies hydride formation. As seen in the above brief review, there remain significant differences in observations of the effects of H on the flow stress of metals. The present paper discusses one aspect of this problem, the effects of H on shear localization, which may account for the discrepancies. Discussion One factor which can account for the differences in observations of the H effects on flow stress is the effect of slip localization on the measured flow stress. Localization of deformation on a limited number of slip planes due to hydrogen charging was observed in stainless steels (19), Ni (20), and high purity AI (16). In the stainless steel-H alloys, the strain localization was accompanied by significant increases in the flow stress, while in the A1-H alloys, significant decreases in the flow stresses was observed. In common with other types of systems, e.g. precipitation hardened alloys, radiation damaged systems, etc., shear localization occurs if the f l o w stress in the region of localization is reduced r~lative to the flow stress of the non-deforming volume. This can occur in several ways as discussed below. Case I. In systems where H reduces the barriers to dislocation motion, an inhomogeneous distribution of H can cause shear localization since the flow stress is lower where the H concentration is greatest. This may occur if H entry is facilitated at slip line intersections with the surface, leading to continuation of slip along active slip bands. If the localization of slip is due to increased dislocation mobility due to H, the tensile flow stress will be decreased. A situation such as this may describe the adsorption of H from gas phases or from liquids, as in charging by chemical or electrochemical charging during deformation, and is consistent with the decreased flow stresses observed by Kimura et al in Fe (12-14) and in plasma charging of Fe (15). Case II. Shear localization can occur when the specimen is hardened by the introduction of hydrogen (due to the formation of hydrides or H clusters), if the initial deformation locally reduces the stress for continued slip, a localization similar to that observed in precipitation hardened systems. This shear localization is probably what occurs in the Ni-H system where the introduction of H forms a surface hydride and the measured flow stress in a tension test increases. Slip localization in homogeneous stainless steel alloys with high hydrogen concentrations is also observed and the tension flow stress is increased. In this case, the localization may be due to the formation of a "hydride phase" at the surface (21) or to the formation of hydrides at dislocation cores. Localization due to removal of H clusters by intersecting dislocations is thought to be the cause of localization in the AI-H system (16), a system in which significant decreases in the flow stress are also observed. While slip localization can occur, independent of any changes in the dislocation mobility (due to shearing of second phases), increases in dislocation mobility due to the presence of solute H will enhance slip localization in all cases. There clearly is no simple correlation between the flow stress measured in a tensile test and the microscopic effects of H on dislocation behavior. The main point is that in the presence of slip localization, a tensile test will not determine the true flow stress of the material. Thus the macroscopic flow stress measured in a tensile test can be increased o1"

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decreased, depending on the magnitude of the hydrogen enhanced dislocation velocity and on the degree of slip localization. In a tensile test, the usual assumption is that the deformation occurs relatively uniformly over the gage length, 1. The tensile machine imposes a deflection rate, i, and the specimen responds by deformation at a strain rate, e, defined as i/1 o or d(ln(1/lo))/dt with the response function of the specimen being the flow stress, zf. The strain rate is related to the dislocation parameters by: = p=bv0z~

(1)

where b is the Burgers vector, Pm is the mobile dislocation density, zf is the resolved shear stress, m is a material dependent exponent having values between 1 and about 30, and v o is a material constant. In the absence of any effect of H on the dislocation properties, the ratio of the flow stress for the specimen undergoing shear localization, %1, to one which is uniformly deforming, "¢u is given by:

'~/~o~/-(~o/~ _

(2)

k/1U

where ll/lu is the fraction of the gage length in which localized shear occurs. The parameter ll/lu can be approximated by the ratio of the slip line density in the specimen exhibiting localized slip and that which exhibits uniform slip. The reason for increased flow stresses due to strain localization is that as the effective gage length is decreased by localization, the effective dislocation density is decreased and the dislocation velocity must be correspondingly increased to match the imposed deflection rate. As indicated in Eqn. 1, this results in an increased flow stress. Values for the increased measured flow stress due to shear localization in the absence of any H related increases in dislocation mobility are given in Table II for a range of ll/l u and m. As seen in Table I, significant increases in the measured flow stress can occur with relatively moderate shear localization. Hence if H causes shear localization and no enhanced dislocation mobility, the macroscopic flow stress will increase; a reflection of the slip localization rather than the effects of H on dislocation behavior. TABLE I. INCREASED FLOW STRESSES DUE TO SHEAR LOCALIZATION IN THE ABSENCE OF HYDROGEN SOFTENING

lu / II

10 102 103 104

"CI/ '~u m=10

m=l

m=5

10 102

1.58 2.51 3.98 6.31

103 104

1.26 1.59 2.00 2.51

m=30

1.08 1.17 1.26 1.36

If H alters the dislocation response to an applied stress, the effects of shear localization on the measured flow stress is somewhat more complex. The measured flow stress, z, can be written:

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j

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.


Since the displacement rate (i) is set by the testing machine, the ratio of the flow stress with and without H is written:

~" -~--

[P~"V°]Y~

(4)

[p~,,Vo,.ly~.

where the subscript H indicates the value in the presence of H. In the presence of both shear localization and hydrogen effects on dislocation mobility, the flow stress ratio can be written:

•,.,.,:r:_.,-ir,,-,]_r,_,lr (,,=,v:),,. ] '< L',::jI..',:..I-Ll, JL(p=.,.,Vo,,.,)"<"

(5)

where "~I-Uis the flow stress in the presence of H and slip localization. The parameters necessary to discuss these equations in detail have not been measured as a function of H concentration. It is evident however, that the flow stress measured in a tensile test can either be increased or decreased by the presence of H depending on the extent of shear localization, ('Cl/'~u), and the effect of H on the dislocation mobility, ('~H/'0. Thus measurement of the effects of H on the macroscopic flow stress alone is not a reliable means of determining dislocation behavior. Measurements of the macroscopic stress strain behavior which show increases in the flow stress due to the introduction of H may result from shear localization, even in the presence of reduced barriers to dislocation motion. The in situ TEM observations suggest that the introduction of H increases pm as well as the velocities of individual dislocations. Increases in dislocation velocity due to H suggests that both mH and Vo,H are increased by the H solutes. From Eqn. 5: macroscopic hardening in a tension test is expected for

11,)

P~Vo

and macroscopic softening in a tensile test is expected for

Il ,)

p,~Vo

(6)

Conclusions 1. The effect of hydrogen on the macroscopic flow stress of metals, as measured by tensile tests, depends both on the effects of H on the dislocation stress velocity relation and on the tendency for H to cause strain localization. 2. Strain localization occurs when the flow stress in the active slip planes is less than the flow stress in the bulk of the specimen. This may result from H enhanced dislocation velocity and inhomogeneous H distributions or from decreases in the barriers to dislocation motion caused by hydrides, etc. when they are cut by dislocations.

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3. Macroscopic hardening, i.e. increases in the tensile flow stress always results from strain localization unles~ softening due to H effects on dislocation behavior are sufficiently large. Thus increases in the tensile flow stress do not necessarily indicate decreased dislocation mobility due to H, but may be due to slip localization caused by H. A decrease in the tensile flow stress does however require that the dislocation mobility be increased by the presence of H. 4. The in situ TEM data suggests that the mobilities of dislocations under stress are increased by the introduction of H into solid solution. This is directly consistent with decreased tensile flow stresses in Fe and with macroscopic softening in the presence of slip localization as observed in high purity A1-H alloys.

5. Increased tensile flow stresses are observed in the presence of slip localization in stainless steel-H and Ni-H alloys despite the TEM observations which show increased dislocation mobilities caused by hydrogen and similar observations are made in steels. This can be accounted for by extensive H induced strain localization. Acknowledgments v

This research was supported the Department of Energy contract DEFG02-91ER45439. References 1. T.C. Lee, D. Dewald, J. A. Eades, I. M. Robertson, H. K. Birnbaum, Rev. of Scientific Instr., 62, 1438 (1991). 2. I.M. Robertson, H. K. Birnbaum, Acta Metall., 34, 353 (1986). 3. T. Tabata, H. K. Birnbaum, Scripta Metall., 18, 231 (1984). 4. D. Shih, I. M. Robertson, H. K. Birnbaum, Acta Metall., 36, 111 (1988). 5. G. Bond, I. M. Robertson, H. K. Birnbaum, Acta Metall., 35, 2289 (1987). 6. I.M. Robertson, H. K. Birnbaum, unpublished, (1987). 7. I.M. Robertson, T. Tabata, W. Wei, F. Heubaum, H. K. Birnbaum, Scripta Metall., 18, 841 (1984). 8. H.K. Birnbaum, P. Sofronis, Materials Science and Eng., in press, (1994). 9. P. Sofronis, H. K. Birnbaum, Jnl. Mech. and Phys. of Sol., submitted. 10. I. M. Bernstein, Scripta Met., 8, 343 (1974). 11. M. Cornet, S. Talbot-Besnard, Japan Inst. Metals, Supl., 21, 5454 (1980). 12. H. Matsui, H. Kimura, S. Moriya, Mat. Sci. and Eng., 40, 207 (1979). 13. S. Moriya, H. Matsui, H. Kimura, Mat. Sci. Eng., 40, 217 (1929). 14. H. Matsui, H. Kimura, A. Kimura, Mat. Sci. Eng., 40, 227 (1979). 15. A. Kimura, H. K. Birnbaum, Scripta Metall., 21, 53-57 (1987). 16. F. Zeides, Ph.D., Univ. of Illinois (1986). 17. E. Sirois, H. K. Birnbanm, Acta Metall., 40, 1377-1385 (1992). 18. A. Cottrell, Phil. Mag., 44, 829 (1953). 19. D. G. Ulmer, C. J. Altstetter, Acta Metall., 28, 1237 (1991). 20. W. A. McLnteer, A. W. Thompson, I. M. Bernstein, Acta Metall., 28, 887 (1980). 21. H. Kimura, H. K. Birnbaum, Acta Metall., 35, 1077 (1987).