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The effect of hydrogen on the mechanical properties of high purity iron I. Softening and hardening of high purity iron by hydrogen charging during tensile deformation
Materials Science and Engineering, 40 (1979) 207 - 216 © Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands
207
The Effect of Hydrogen on the Mechanical Properties of High Purity Iron
I. Softening and Hardening of High Purity Iron b y Hydrogen Charging during Tensile Deformation H. MATSUI and H. KIMURA Research Institute for Iron, Steel and Other Metals, Tohoku University, Sendai (Japan) S. MORIYA* Graduate School of Engineering, Tohoku University, Sendai (Japan) (Received March 31, 1979)
SUMMARY Tensile tests were performed on high purity iron during hydrogen charging. The flow stress was increased by hydrogen charging below 190 K and was decreased reversibly above 190 K. Hydrogen charging caused no, or only very few, blisters on the surface o f specimens with a residual resistivity ratio ( R R R H ) larger than 3600 but an appreciable number o f blisters on less pure specimens (RRRH ~ 1800). The less pure specimens were hardened rather than softened by hydrogen charging at room temperature. Several mechanisms for the softening are discussed, and it is concluded that the softening and hardening observed in the high purity specimens are the inherent effect of hydrogen on the mechanical properties o f iron. The softening and hardening are interpreted as the results o f hydrogen-dislocation interactions; it is proposed that hydrogen atoms trapped at the core of a screw dislocation increase its mobility and that hydrogen atoms hinder the motion o f edge dislocations at sufficiently low temperatures.
1. INTRODUCTION It is well known that hydrogen embrittles ferrous alloys. Accordingly, great efforts have been made to investigate the effect of hydrogen on the mechanical properties of iron and
*Present address: Kinuura Works, Nippon Metal Industry Co. Ltd., Japan.
steels. Nevertheless, there is serious confusion in published reports, especially on iron. The most important controversy is over the question: "Does hydrogen harden or soften iron?" The majority of reports [1 - 3] treating " p u r e " iron conclude that hydrogen hardens iron. Reports [4] stating that hydrogen softens iron are in a minority. The most important cause of this controversy may be the small solid solubility of hydrogen in iron [5] as compared with the a m o u n t of impurity contained in the " p u r e " iron used in most of the researches. The solid solubility and the diffusivity of hydrogen in iron are greatly affected by impurities in iron. Thus, it is necessary to use very pure specimens in which the concentration of impurities is less than that of hydrogen in order to reveal the true effect of hydrogen on the mechanical properties of iron. It has already been reported briefly by the present authors [6] that high purity iron is softened b y electrolytic hydrogen charging during tensile testing. Several possible causes of the softening were discussed, and it was suggested that the interaction of hydrogen with screw dislocations is the most likely cause of the softening [6, 7]. In the present research the experimental conditions were wider than in our previous research; for example, the testing temperature was lowered to 170 K and detailed transmission electron microscopy was performed. Consequently, the effects of hydrogen charging on the mechanical properties of high purity iron were revealed much more comprehensively than in the previous research [6] and hence a more thorough discussion of the cause of softening has become possible.
208 2. EXPERIMENTAL
2.1. Specimens The J o h n s o n - M a t t h e y (JM) iron rods of 5 mm diameter were floating zone refined by electron beam under an ultrahigh vacuum of 1.3 × 10 -6 - 1.3 × 10 -7 Pa for more than 12 passes. The purified rods were cold swaged to 2.2 mm diameter and were then cold drawn to 0.5 mm diameter without intermediate annealing. The drawn wire was hydrogen treated for 1 h at 1200 K in dry hydrogen, for 24 h at 1070 K in wet hydrogen and finally for 20 h at 1070 K in dry hydrogen. The hydrogen gas was purified by a palladium-permeation purifier. The surface layer contaminated with silicon during dry hydrogen annealing [ 8] was polished off and the resulting specimen was usually 0.4 m m in diameter. The purity of the specimens was estimated using the residual resistivity ratio (RRRH), with the resistivity at 4.2 K being measured in a magnetic field of about (2/u) × 105 A m - 1 . For a more detailed description of the purification see ref. 9. The specimens were classified into three groups according to purity, i.e. using RRRH : grade A with RRRH ~ 5000, grade B with RRRH ~ 3600 and grade C with RRRH 1800. The grade C specimens were of hydrogen-treated JM iron that had n o t been zone refined. The impurities in the present specimens were not analyzed. However, a mass spectroscopic analysis was performed on a high purity iron produced in the same way [10], and the total impurity concentration was found to be about 10 at.ppm (atomic parts per million). The carbon in solution was estimated to be about 2 at.ppm [9]. Most of the specimens for tensile testing were 0.4 mm in diameter and 10 - 20 m m in gauge length. The grain sizes were about 1.0, 0.3 and 0.1 mm respectively for the three grades of purity; the purer specimens had the larger grain size. The grade A and B specimens had a bamboo-type structure. The specimens for hydrogen analysis were 2 mm in diameter. The specimens for transmission electron microscopy (TEM) were cut from cold-rolled foils with dimensions of 0.15 m m × 2 m m after purification by hydrogen annealing as mentioned above. They were subsequently polished to a thickness of 0.12 m m to remove the silicon-contaminated layer.
2.2. Hydrogen charging Hydrogen was charged electrolytically mostly in a solution of 0.1 N C H s O H - H 2 0 H 2 S O 4 containing a small a m o u n t of N a A s O 2 . Other electrolytes were also used for comparison, e.g. dilute aqueous H 2 S O 4 solution, NaOH solution and H2SO 4 diluted with C2H5OH , all containing NaAsO2. Softening due to hydrogen charging was observed with all these solutions the most softening being obtained with the C H 3 O H - H 2 0 H 2 S O 4 solution. The difference in the degree of softening is considered to be due to the difference in the effective hydrogen pressure. The temperature rise in the specimen due to Joule heating during charging was measured to examine its effect on the flow stress, since the flow stress in iron depends quite strongly on temperature. The specimen temperature was measured with a copper-constantan thermocouple in contact with the specimen across a thin layer of insulating lacquer. The temperature rise was less than 0.5 K even at a total charging current of 0.5 A, which is about three orders of magnitude larger than that in the present experiment, i.e. 0.25 - 0.5 mA. Thus the effect of Joule heating is negligible. The charging current density was 20 A m-2 unless otherwise stated. This current density corresponds to the total current of 0.25 mA and it gave the maximum amount of softening at all temperatures examined. The a m o u n t of hydrogen in specimens with RRRH = 3600 charged at 250 K was measured with an argon gas carrier h o t extraction method. It took about 10 min to start the extraction. A very considerable a m o u n t of hydrogen may have escaped out of these specimens before the analysis [11]. Hence the results of the analysis should be regarded as greatly underestimated. They range from 15 to 30 at.ppm.
2. 3. Mechanical testing Tensile experiments were carried out with an Instron-type tensile testing machine. An electrolytic cell surrounded with heater wire was attached to the tensile jig. The cell was in thermal contact with liquid nitrogen through a copper ribbon attached to its b o t t o m . The bath temperature was controlled by regulating the current through the heater; the precision of the temperature control was +0.1 K.
209
3.2.1. Above 190 K
3. R E S U L T S
3.1. Blistering by hydrogen charging Hydrogen charging sometimes damages the specimen by producing internal cracks and surface blistering, and the damage causes softening in mild steel and electrolytic iron specimens [2, 12]. Hence the surface of the specimens was carefully examined with an optical microscope. No blisters were found on the surface of the grade A specimens (RRRH 5000), a few blisters were found in some of the grade B specimens ( R R R H ~ 3600), and m a n y blisters in the grade C specimens (RRRH ~ 1800).
3.2. The effect o f deformation temperature Figure 1 shows several tensile stress-tensile strain curves for the grade A and B specimens tested at a strain rate of 8.3 X 10 -5 s-1 at various temperatures. Tensile tests were started without hydrogen charging. The downward and upward arrows show the charging current being switched on and off respectively. The behavior of the flow stress may be classified into two types, i.e. above and below 190 K.
The flow stress decreases promptly when the charging current is switched on. It decreases rapidly at first and eventually reaches a steady level. Occasionally, the flow stress decreased in two steps around 250 K. However, the two-step decrease was not easily reproducible so it will not be discussed here. The a m o u n t of stress decrease A o is as large as about 50% of the flow stress level immediately before charging. The flow stress begins to increase on switching off the current and recovers to roughly the same stress level as before charging. The on-and-off cycle of the charging current can be repeated reversibly several times especially at high temperatures, e.g. 294 K, before fracture occurs. The fracture is exclusively intergranular as observed by scanning electron microscopy.
3.2.2. Below 190 K On application of the charging current the flow stress decreases only temporarily and then increases steeply (Fig. 1). The maximum stress is u n k n o w n since fracture occurs pre300,
250 RRRH ; 3 6 0 0 - 5 2 0 0 ~=
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200
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,
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,
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,
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Fig. 1. T h e e f f e c t o f hydrogen charging on the flow stress o f h i g h p u r i t y i r o n at various l o w t e m p e r a t u r e s .
Tensile Stroin
Fig. 2. T h e e f f e c t o f hydrogen charging on the flow stress o f high purity iron deformed a t 170 K at three d i f f e r e n t s t r a i n rates.
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maturely. Figure 2 shows the stress-strain curves at 170 K at three different strain rates but with the same charging current density. Fracture also occurs prematurely here but the fracture stress rises to 290 MPa at the two strain rates larger than 8.3 × 10 - 5 s -z . Hence we consider that at 180 and 170 K hydrogen charging causes hardening and that the softening immediately after application of the current is only transient. It is not known whether hardening would be observed at 190 K if fracture did not occur. Figure 3 summarizes the temperature effect for the grade A and B (RRRH = 3 5 0 0 - 5200) specimens. The upper curve represents the flow stress immediately before application of the charging current and the lower curve is the minimum flow stress during charging. The amount of stress decrease A o (the difference between the two curves) is largest near 200 K and decreases at higher temperatures. The solid symbols in the figure represent the fracture stress. The effect of hydrogen charging
changes drastically from softening to hardening as the test temperature decreases below about 190 K. It should be noted here that the temperature dependence of the flow stress above 200 K is markedly decreased by hydrogen charging.
3.3. The effect of purity Figure 4 shows the effect of hydrogen charging on the stress-strain curves of the grade C specimens (RRRH ~ 1800). At low temperatures the effect is similar to that in higher purity specimens, although the amount of stress decrease IA o l is smaller for the grade C specimens. At higher temperatures a small hump appears immediately after application of the current; the flow stress then decreases rather gradually. The hump grows as the temperature is raised and eventually at 273 K no decrease in the flow stress is observed. Thus near room temperature only hardening is observed. Softening still occurs in these specimens if a large current is applied, e.g.
300 = 8.3x RRR. o 0.1% v 250
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Fig. 3. The amount of softening due to hydrogen charging at various temperatures. Hardening and fracture take place below 190 K.
Strain
:5
('/.)
Fig. 4. The effect of hydrogen charging on the flow stress of low purity iron with R R R H ~ 1 5 0 0 .
211
80 A m-2 as shown by the curve at the bott o m of the figure.
3.4. Electron microscopy Foil specimens (0.12 mm thick, 2 mm wide, 20 mm long) of grade B were tensile tested mainly at room temperature and at 200 K with and without hydrogen charging. TEM specimens were prepared by jet electropolishing from the foil tensile specimens and were examined in an electron microscope operating at 100 kV. Figure 5 shows a pair of electron micrographs of specimens deformed at 200 K. Figure 5(a) shows an uncharged specimen (e ~ 3.8%) and Fig. 5(b) a charged specimen (e ~ 3.0%, ic = 8 A m-2) *. The dislocation density is considerably increased in the hydrogen-charged specimen. By surveying several TEM specimens over a wide area it was found that in uncharged specimens the dislocation density varies from 3 × 10 s to 6 × 109 cm-2 from grain to grain whereas in charged specimens it varies in the range (3 - 6) X 109 cm -2 . The average dislocation density is larger by a factor of 5 - 10 in the charged specimens than in the uncharged specimens. The specimens deformed at room temperature with and without charging show no appreciable difference in dislocation density, which is in the range (1 - 5) × 108 cm -2 . Accurate determination of the dislocation character was impossible because none of the grains were found to be of a suitable orientation for a tilting experiment. No significant difference was found in the general features of the dislocation configuration between uncharged and charged specimens.
I10
g -- 2 0 0 I Fm
ITO
(a)
g = 200 4. D I S C U S S I O N
4.1. The mechanism o f softening The most remakable effect of hydrogen charging is the softening in the high purity specimens {grade A and B) at all temperatures above 190 K and in the less pure grade C specimens below room temperature. Several possible mechanisms for the softening have
* F o r a foil specimen the current density o f 8 A -2 m gives an a m o u n t o f softening similar t o t h a t in wire specimens. The effect o f specimen size on the s o f t e n i n g is r e p o r t e d in a separate paper [ 1 3 ] .
I Fm
(b) t
I
Fig. 5. T r a n s m i s s i o n e l e c t r o n m i c r o g r a p h s o f h i g h p u r i t y i r o n s p e c i m e n s d e f o r m e d at 2 0 0 K: ( a ) ' w i t h o u t h y d r o g e n charging, e ~ 3.8%; ( b ) w i t h h y d r o g e n charging, i c = 8 A m - 2 , e ~ 3.0%.
already been discussed to show that the softening is probably due to an interaction between hydrogen and screw dislocations [6]. However, the previous discussion was not necessarily complete, and a further and more
212 complete discussion is given below on the basis of the present new results.
4.1.1. Damage caused by hydrogen charging Softening may occur if the specimen crosssectional area is reduced b y the formation of cracks as a result of hydrogen gas precipitation. However, this mechanism is ruled out because the softening is reversible on switching the current on and off. Reversible softening by hydrogen charging is observed in mild steel and electrolytic iron when the charging current density is large and/or the specimen is thin enough [3, 1 2 ] . The softening is considered to be due to dislocation generation at damage, e.g. at blisters and internal cracks [2, 3, 1 2 ] . Hydrogen gas pressure built up in the damage during charging m a y enhance the dislocation generation. There is evidence to rule out damage as the cause of the softening in high purity iron. In the present research, blisters are formed in the grade C specimens which show hardening, not softening, at room temperature. The grade C specimens show softening at lower temperatures but the amount/x o of softening is smaller than in the grade B and A specimens in which very few or no blisters are formed. It is reasonable to assume that the number, if any, of internal cracks is also smaller in the grade A and B specimens than in the grade C specimens. Moreover, it might be difficult to explain the hardening observed below 190 K using the damage hypothesis.
4.1.2. Surface effects Some electrochemical reactions occurring during charging could cause softening, as in the case of the Rehbinder effect [ 1 4 ] . However, this mechanism is ruled out because very similar softening occurs in specimens {grade A) charged at room temperature and then quenched to and tested at low temperatures. In this experiment specimens were transferred from the electrolytic cell into a nonelectrolytic bath at low temperatures. Details of this series of experiments, which will be referred to as the "after-charging experiment", will be reported in a separate paper [ 1 5 ] . The softening could be caused by a steep solute concentration gradient. This t y p e of softening has been suggested by Sethi and Gibala [16] for the softening of surface oxide films. Their mechanism assumes that a large stress
produced by a steep concentration gradient at the surface nucleates edge dislocations, which are mobile under a much lower stress than screw dislocations. In the present experimental conditions the fugacity of hydrogen on the surface is of the order of 106 - 101° Pa [17] and the corresponding lattice solubility of hydrogen at room temperature is estimated to be 4 - 400 at.ppm by Gonzalez' measurement [5]. This solubility may give a hydrogen concentration gradient of 400 at.ppm per atomic distance at the most. In contrast, the hydrogen concentration gradient required to nucleate edge dislocations is estimated to be about 5 at.% per atomic distance [16] for a partial molar volume of hydrogen in iron of 2.0 cm ~ (mol H) -1 [18]. Thus this mechanism cannot be the cause of the present softening. Moreover, in the "after-charging experiment" mentioned above, the concentration gradient, if any, has probably disappeared before the tensile test is started because the fast diffusing hydrogen escapes from the surface layer immediately when the current is switched off. Nevertheless, softening is observed.
4.1.3. Bulk effects The above discussion together with those in previous papers [6, 7, 19] enables us to conclude that the softening is a bulk effect. In the following, several mechanisms will be discussed and it will be shown that the softening is most probably due to a hydrogenscrew dislocation interaction; hydrogen trapped at the core of screw dislocations increases their mobility.
4.1.3.1. The scavenging effect. Hydrogen might scavenge impurities already present in iron and remove the hardening effect of these impurities. This mechanism is readily discarded since, as stated in Section 3.3, the purer specimens show more pronounced softening. 4.1.3.2. Dislocation punching from precipitates. In a Va metal-hydrogen system hydride precipitates punch out edge dislocations, which are much more mobile than the screws, and cause softening [20]. It might be possible for hydride or hydride-like precipitates to exist in iron under an extraordinarily high effective pressure of hydrogen pro-
213 duced by charging. If such precipitates do exist, they may cause softening. Baranowski has suggested that iron hydride will exist, if at all, only under an extremely high hydrogen fugacity which can be estimated to be larger than 1021 Pa [21]. The effective hydrogen fugacity in the present experiments is considered to be less than 101° Pa. Hence the hydride hypothesis is unlikely. The result of the "after-charging experiment" is also unfavorable for the hydride hypothesis, as reported separately [15].
4.1.3.3. Increased dislocation density. Since TEM observation shows that the dislocation density in hydrogen-charged and deformed specimens is larger than in specimens without charging, we might consider that the softening is due to the increased dislocation density. As will be discussed below, the observed increase in dislocation density is not sufficient to explain the present softening. TEM observation shows that most of the dislocations are screws in both hydrogen-charged and uncharged specimens. The strain rate sensitivity of the flow stress at 200 K without charging has been found to be 20 - 25 MPa for a change in the strain rate by a factor of 10 [22]. The dislocation density is found to increase by a factor of 5 - 10 with hydrogen charging. This increase may cause a flow stress decrease of 10 - 25 MPa at 200 K if the strain rate is kept unchanged. The observed decrease in stress due to hydrogen charging at 200 K is about 80 MPa, which is much larger than the estimated value. To cause the observed softening the dislocation density would have to, increase by a factor of 103 - 104 . However, if the dislocation density were increased by this a m o u n t , it would cause an appreciable increase in the internal stress and hence hardening. Thus it is clear that the observed increase in mobile dislocation density is certainly insufficient to explain the observed softening. 4.1.3.4. The interaction of hydrogen with screw dislocations. The above discussion rules out all mechanisms for softening proposed so far except one, i.e. the interaction of hydrogen with screw dislocations. Solution softening and hardening have been observed in the iron-carbon [23] and iron-nitrogen systems [24] as well as in irradiated iron [25].
The softening and hardening have been discussed in terms of the interaction of screw dislocations with solute atoms and point defects [26]. The temperature range of the softening and hardening due to carbon and nitrogen is very similar to that for the present softening and hardening due to hydrogen. The similarity, however, does not mean that the same mechanism is operating for hydrogen as for carbon and nitrogen, since the behavior of hydrogen in iron in this temperature range is very different from that of carbon and nitrogen. Carbon and nitrogen are practically immobile below room temperature where they cause softening. Radiation softening is observed at temperatures where radiationinduced point defects are considered immobile. The hydrogen diffusivity, in contrast, is quite large (i.e. 6.3 × 1 0 - 6 c m 2 s - 1 [27] ) even at 200 K and hydrogen is considered to move together with dislocations. Hence theories of softening in which immobile solute atoms dispersed in the matrix enhance the double kink nucleation on screw dislocations by an elastic interaction cannot be applied to the present case. Since hydrogen is quite mobile it should be assumed to enter the dislocation core and m o d i f y the core structure to increase the kink nucleation rate. Hydrogen could be immobilized by forming hydrogen pairs, h y d r o g e n - i m p u r i t y pairs etc. which could act as stress centers to facilitate kink pair nucleation. This is unlikely, however. A resistivity study of the diffusion of hydrogen in high purity iron (RRRH >~ 3000) showed that excess hydrogen in the matrix escapes from the surface below 170 K and gave no evidence for the presence of a detectable number of immobile pairs in the matrix above 200 K [11]. Nevertheless, the "after-charging e x p e r i m e n t " reveals softening, as described in ref. 15. This observation rules out the hypothesis that immobilized hydrogen exists in the matrix and causes the softening. Moreover, if h y d r o g e n - i m p u r i t y pairs were the cause of softening, more softening should have been found in the less pure specimens; in fact the opposite was observed.
4.2. Proposed mechanisms for the softening and hardening If hydrogen atoms bound to the screw dislocation core are assumed to increase the double kink nucleation rate and hence the
214 screw dislocation mobility, the softening is readily explained. The hardening below 190 K is explained in terms of the hydrogen-edge dislocation interaction. Kinks on a screw dislocation move along it much more quickly than the screw dislocation itself. If the diffusivity of hydrogen atoms along a screw dislocation is sufficiently reduced below 190 K, hydrogen atoms may act as dragging points to the kinks. Since the Peierls potential of an edge dislocation is much smaller than that of a screw dislocation, hydrogen atoms may pin an edge dislocation rather than increase its mobility by reducing the Peierls stress. Thus below 190 K kinks are easily formed with the help of a hydrogen atmosphere on a screw dislocation but their sideways motion is hindered by hydrogen; hence hardening results. There is evidence that hydrogen pins the kinks (see ref. 15). The present experimental results, other than those discussed above, will be explained in terms of the proposed mechanisms. The transient softening observed at and below 180 K may be due to competitive operation of the softening and hardening mechanisms. For example, the following process may be suggested. The softening occurs at a relatively small concentration of hydrogen atmosphere because the presence of only a few sites of easy kink nucleation facilitates the motion of a screw dislocation as a whole. The hardening requires more hydrogen, since the easily nucleated kinks must be pinned before they move a considerable distance along the screw dislocations. According to this mechanism, softening will be observed first on application of the charging current, and hardening will take over when a sufficient amount of hydrogen has entered into the specimen. The result that more softening is observed at lower temperatures may partly be due to larger hydrogen concentrations in the atmosphere at lower temperatures. However, this effect would not be very appreciable. The atmosphere concentration is proportional to the concentration in the matrix, which is related by Sieverts' law to the effective pressure surrounding the specimen, at least at temperatures not t o o far below room temperature. Since the heat of solution of hydrogen in iron is negative, as the temperature is lowered, the matrix concentration is also
lowered. Consequently, the hydrogen concentration in the core will not increase appreciably as the temperature decreases. It is more likely that the temperature dependence of the softening is due to the reduction of the activation energy for the screw dislocation motion. In the above discussion it is assumed that the electrolytic reaction and the effective hydrogen pressure on the surface at a constant charging current density ic are independent of the temperature. This assumption is considered to be reasonable for two reasons: (a) similar A o versus i¢ relations are obtained at various temperatures; (b) a yield stress versus temperature relation similar to that in Fig. 3 is obtained in the "after-charging experiment". 4.3. The effect o f p u r i t y The concentration of dissolved impurity atoms is estimated to be about 10 at.ppm in the grade A and B specimens. The hydrogen concentration in the matrix during charging is estimated to be several tens of atomic parts per million [11] and it exceeds the impurity concentration. Hence the softening and hardening in these high purity specimens are considered to be the effect of hydrogen itself. The impurity concentration in the grade C specimens may be comparable with the hydrogen concentration, and it is reasonable to consider the hardening in these specimens at and above 273 K to be due to interactions of hydrogen and impurities with dislocations. As the temperature is lowered, softening, for which the mechanism is the same as for the high purity specimen, becomes predominant. The grade C specimens show softening at 297 K with increasing charging current density. This is probably due to the hydrogenenhanced screw dislocation motion, as in the grade A and B specimens. Since the concentration of hydrogen is larger for the larger current density, the hydrogen concentration charged with 80 A m -2 (Fig. 5) may be considered to exceed a critical concentration determined by impurities, and hence softening results. Blisters are formed in this charging condition (80 A m -2) but they will n o t cause appreciable softening. It should be mentioned here that a current density larger than 300 A m - 2 is necessary to cause damage-induced softening in an electrolytic iron with a similar electrolytic solution [ 1 2 ] . The degree of damage is much larger in an
215
impure iron than in a high purity iron. Hence damage caused by charging with 80 A m - 2 produces hardly any softening. 4.4. The fast response of the flow stress change The flow stress responds quickly to switching on of the charging current because of transportation of hydrogen by moving dislocations. Hydrogen also diffuses in towards the specimen interior by bulk diffusion. Although it is n o t known at present what fraction of hydrogen is transported by moving dislocations until the flow stress attains the steady level, the dislocation transportation should be dominant in the initial period of flow stress decrease. 4.5. The binding energy between a hydrogen atom and the core of a screw dislocation In the above discussion an appreciable binding energy is assumed to exist between hydrogen and the screw dislocation core. The binding energy between a hydrogen atom and an edge dislocation has been estimated to be a b o u t 0.3 eV [ 2 8 ] . No measurement has been performed, however, for screw dislocations. Chou has estimated on the basis of an elasticity theory that the binding energy of a solute atom with screw dislocations is about one-sixth of that with edge dislocations in iron [ 2 9 ] . Since the difference in the binding energy with the dislocation core will not be as great between edge and screw dislocations as the difference estimated from elasticity, the binding energy between a hydrogen atom and a screw dislocation core may be 0.1 - 0.3 eV. If hydrogen pairs are formed in the dislocation core (not in the matrix), the binding energy with dislocations will be increased. A binding energy of 0.2 eV is enough to form an appreciable concentration of hydrogen atmosphere at and around the screw dislocation core below room temperature. 5. CONCLUSIONS
Tensile tests were performed on high purity iron during hydrogen charging. The flow stress is increased by hydrogen charging below 190 K and decreased reversibly above 190 K. Hydrogen charging causes no, or only very few, blisters on the surface of specimens with RRRH > 3600, but an appreciable number of blisters on less pure specimens (RRRH
1800}. The less pure specimens were hardened rather than softened by hydrogen charging at room temperature. Several mechanisms for the hydrogeninduced softening, including damage due to charging, have been discussed, and it is concluded that the softening in high purity iron is due to an interaction between hydrogen and screw dislocations. The hardening below 190 K is due to an interaction between hydrogen and edge dislocations, including kinks on screw dislocations. The softening and hardening are the inherent effect of hydrogen on the mechanical properties of pure iron. It is proposed that hydrogen atoms at the core of screw dislocations m o d i f y their Peierls potential and thus increase their mobility. ACKNOWLEDGMENTS
This research was partly supported financially by the Ishihara-Asada Fund of the Iron and Steel Institute of Japan (S.M.) and by the Matsunaga Science Foundation (H.M.). The authors are grateful for this support. The authors thank Dr. S. Takaki for his cooperation in preparing the high purity iron specimens and Mr. Hosoya for his hydrogen analysis. REFERENCES 1 A. M. Adair, Trans. Metall. Soc. AIME, 236 (1966) 1613. S. Matsuyama, in Mechanism o f Delayed Failure o f Steels due to Hydrogen, Iron and Steel Institute of Japan, Tokyo, 1975, p. 113. Y. Tobe and W. R. Tyson, Scr. Metall., 11 (1977) 849. 2 S. Asano and R. Otsuka, Scr. Metall., 10 (1976) 1015. 3 S. Asano and R. Otsuka, Scr. MetaU., 12 (1978) 287. 4 C. D. Beacham, Metall. Trans., 3 (1972) 437. I. M. Bernstein, Scr. Metall., 8 (1974) 343. 5 O. D. Gonzalez, Trans. Metall. Soc. AIME, 245 (1969) 607. 6 H. Matsui, S. Moriya and H. Kimura, Proc. 4th Int. Conf. on Strength o f Metals and Alloys, Nancy, France, 1976, p. 291. 7 H. Kimura, H. Matsui and S. Moriya, Scr. Metall., 11 (1977) 473. 8 S. Takaki and H. Kimura, Scr. Metall., 10 (1976) 701. 9 S. Takaki and H. Kimura, Scr. Metall., 10 (1976) 1095.
216
10 S. Takaki, personal communication, 1976. 11 S. Moriya and H. Kimura, Mater. Sci. Eng., 32 (1978) 71. 12 H. Hagi, Y. Hayashi and N. Ohtani, J. Jpn. Inst. Met., 42 (1978) 124. 13 H. Matsui, A. Kimura and H. Kimura, Mater. Sci. Eng., 40 (1979) 227. 14 P.A. Rehbinder, Nature (London), 159 (1947) 866. 15 S. Moriya, H. Matsui and H. Kimura, Mater. Sci, Eng., 40 {1979) 217. 16 V. K. Sethi and R. Gibala, Acta Metall., 25 (1977) 321. 17 A . J . Kumnik and H. H. Johnson, Metall. Trans. A,6 (1975) 1087. 18 R. A. Oriani, Trans. Metall. Soc. AIME, 236 (1966) 1368. 19 H. Kimura and H. Matsui, Scr. Metall., 13 (1979) 221. 20 R. Fournier and R. Gibala, Proc. 2nd Int. Congr. on Hydrogen in Metals, 1977, Pergamon, Oxford, Paper 1E6.
21 B. Baranowski, Ber. Bunsenges. Phys. Chem., 76 (1972) 714. 22 H. Matsui, S. Moriya, S. Takaki and H. Kimura, Trans. Jpn. Inst. Met., 19 (1978) 163. 23 R. B. Roy and W. Stany, Z. Metallkd., 64 (1973) 511. D. J. Quesnell, A. Sato and M. Meshii, Mater. Sci. Eng., 18 (1975) 199. 24 Y. Nakada and A. S. Keh, Acta Metall., 16 (1968) 903. 25 A. Sato, T. Mifune and M. Meshii, Phys. Status Solidi A, 18 (1973) 699. 26 A. Sato and M. Meshii, Acta Metall., 21 (1973) 753. 27 R. A. Oriani, in R. W. Staehle, A. J. F o r t y and D. van Rooyen (eds.), Fundamental Aspects o f Stress Corrosion Cracking, National Association of Corrosion Engineers, Houston, Texas, 1969, p. 32. 28 R. Gibala, Trans. Metall. Soc. AIME, 239 (1967) 574. 29 Y. T. Chou, Acta Metall., 13 (1965) 251.
207
The Effect of Hydrogen on the Mechanical Properties of High Purity Iron
I. Softening and Hardening of High Purity Iron b y Hydrogen Charging during Tensile Deformation H. MATSUI and H. KIMURA Research Institute for Iron, Steel and Other Metals, Tohoku University, Sendai (Japan) S. MORIYA* Graduate School of Engineering, Tohoku University, Sendai (Japan) (Received March 31, 1979)
SUMMARY Tensile tests were performed on high purity iron during hydrogen charging. The flow stress was increased by hydrogen charging below 190 K and was decreased reversibly above 190 K. Hydrogen charging caused no, or only very few, blisters on the surface o f specimens with a residual resistivity ratio ( R R R H ) larger than 3600 but an appreciable number o f blisters on less pure specimens (RRRH ~ 1800). The less pure specimens were hardened rather than softened by hydrogen charging at room temperature. Several mechanisms for the softening are discussed, and it is concluded that the softening and hardening observed in the high purity specimens are the inherent effect of hydrogen on the mechanical properties o f iron. The softening and hardening are interpreted as the results o f hydrogen-dislocation interactions; it is proposed that hydrogen atoms trapped at the core of a screw dislocation increase its mobility and that hydrogen atoms hinder the motion o f edge dislocations at sufficiently low temperatures.
1. INTRODUCTION It is well known that hydrogen embrittles ferrous alloys. Accordingly, great efforts have been made to investigate the effect of hydrogen on the mechanical properties of iron and
*Present address: Kinuura Works, Nippon Metal Industry Co. Ltd., Japan.
steels. Nevertheless, there is serious confusion in published reports, especially on iron. The most important controversy is over the question: "Does hydrogen harden or soften iron?" The majority of reports [1 - 3] treating " p u r e " iron conclude that hydrogen hardens iron. Reports [4] stating that hydrogen softens iron are in a minority. The most important cause of this controversy may be the small solid solubility of hydrogen in iron [5] as compared with the a m o u n t of impurity contained in the " p u r e " iron used in most of the researches. The solid solubility and the diffusivity of hydrogen in iron are greatly affected by impurities in iron. Thus, it is necessary to use very pure specimens in which the concentration of impurities is less than that of hydrogen in order to reveal the true effect of hydrogen on the mechanical properties of iron. It has already been reported briefly by the present authors [6] that high purity iron is softened b y electrolytic hydrogen charging during tensile testing. Several possible causes of the softening were discussed, and it was suggested that the interaction of hydrogen with screw dislocations is the most likely cause of the softening [6, 7]. In the present research the experimental conditions were wider than in our previous research; for example, the testing temperature was lowered to 170 K and detailed transmission electron microscopy was performed. Consequently, the effects of hydrogen charging on the mechanical properties of high purity iron were revealed much more comprehensively than in the previous research [6] and hence a more thorough discussion of the cause of softening has become possible.
208 2. EXPERIMENTAL
2.1. Specimens The J o h n s o n - M a t t h e y (JM) iron rods of 5 mm diameter were floating zone refined by electron beam under an ultrahigh vacuum of 1.3 × 10 -6 - 1.3 × 10 -7 Pa for more than 12 passes. The purified rods were cold swaged to 2.2 mm diameter and were then cold drawn to 0.5 mm diameter without intermediate annealing. The drawn wire was hydrogen treated for 1 h at 1200 K in dry hydrogen, for 24 h at 1070 K in wet hydrogen and finally for 20 h at 1070 K in dry hydrogen. The hydrogen gas was purified by a palladium-permeation purifier. The surface layer contaminated with silicon during dry hydrogen annealing [ 8] was polished off and the resulting specimen was usually 0.4 m m in diameter. The purity of the specimens was estimated using the residual resistivity ratio (RRRH), with the resistivity at 4.2 K being measured in a magnetic field of about (2/u) × 105 A m - 1 . For a more detailed description of the purification see ref. 9. The specimens were classified into three groups according to purity, i.e. using RRRH : grade A with RRRH ~ 5000, grade B with RRRH ~ 3600 and grade C with RRRH 1800. The grade C specimens were of hydrogen-treated JM iron that had n o t been zone refined. The impurities in the present specimens were not analyzed. However, a mass spectroscopic analysis was performed on a high purity iron produced in the same way [10], and the total impurity concentration was found to be about 10 at.ppm (atomic parts per million). The carbon in solution was estimated to be about 2 at.ppm [9]. Most of the specimens for tensile testing were 0.4 mm in diameter and 10 - 20 m m in gauge length. The grain sizes were about 1.0, 0.3 and 0.1 mm respectively for the three grades of purity; the purer specimens had the larger grain size. The grade A and B specimens had a bamboo-type structure. The specimens for hydrogen analysis were 2 mm in diameter. The specimens for transmission electron microscopy (TEM) were cut from cold-rolled foils with dimensions of 0.15 m m × 2 m m after purification by hydrogen annealing as mentioned above. They were subsequently polished to a thickness of 0.12 m m to remove the silicon-contaminated layer.
2.2. Hydrogen charging Hydrogen was charged electrolytically mostly in a solution of 0.1 N C H s O H - H 2 0 H 2 S O 4 containing a small a m o u n t of N a A s O 2 . Other electrolytes were also used for comparison, e.g. dilute aqueous H 2 S O 4 solution, NaOH solution and H2SO 4 diluted with C2H5OH , all containing NaAsO2. Softening due to hydrogen charging was observed with all these solutions the most softening being obtained with the C H 3 O H - H 2 0 H 2 S O 4 solution. The difference in the degree of softening is considered to be due to the difference in the effective hydrogen pressure. The temperature rise in the specimen due to Joule heating during charging was measured to examine its effect on the flow stress, since the flow stress in iron depends quite strongly on temperature. The specimen temperature was measured with a copper-constantan thermocouple in contact with the specimen across a thin layer of insulating lacquer. The temperature rise was less than 0.5 K even at a total charging current of 0.5 A, which is about three orders of magnitude larger than that in the present experiment, i.e. 0.25 - 0.5 mA. Thus the effect of Joule heating is negligible. The charging current density was 20 A m-2 unless otherwise stated. This current density corresponds to the total current of 0.25 mA and it gave the maximum amount of softening at all temperatures examined. The a m o u n t of hydrogen in specimens with RRRH = 3600 charged at 250 K was measured with an argon gas carrier h o t extraction method. It took about 10 min to start the extraction. A very considerable a m o u n t of hydrogen may have escaped out of these specimens before the analysis [11]. Hence the results of the analysis should be regarded as greatly underestimated. They range from 15 to 30 at.ppm.
2. 3. Mechanical testing Tensile experiments were carried out with an Instron-type tensile testing machine. An electrolytic cell surrounded with heater wire was attached to the tensile jig. The cell was in thermal contact with liquid nitrogen through a copper ribbon attached to its b o t t o m . The bath temperature was controlled by regulating the current through the heater; the precision of the temperature control was +0.1 K.
209
3.2.1. Above 190 K
3. R E S U L T S
3.1. Blistering by hydrogen charging Hydrogen charging sometimes damages the specimen by producing internal cracks and surface blistering, and the damage causes softening in mild steel and electrolytic iron specimens [2, 12]. Hence the surface of the specimens was carefully examined with an optical microscope. No blisters were found on the surface of the grade A specimens (RRRH 5000), a few blisters were found in some of the grade B specimens ( R R R H ~ 3600), and m a n y blisters in the grade C specimens (RRRH ~ 1800).
3.2. The effect o f deformation temperature Figure 1 shows several tensile stress-tensile strain curves for the grade A and B specimens tested at a strain rate of 8.3 X 10 -5 s-1 at various temperatures. Tensile tests were started without hydrogen charging. The downward and upward arrows show the charging current being switched on and off respectively. The behavior of the flow stress may be classified into two types, i.e. above and below 190 K.
The flow stress decreases promptly when the charging current is switched on. It decreases rapidly at first and eventually reaches a steady level. Occasionally, the flow stress decreased in two steps around 250 K. However, the two-step decrease was not easily reproducible so it will not be discussed here. The a m o u n t of stress decrease A o is as large as about 50% of the flow stress level immediately before charging. The flow stress begins to increase on switching off the current and recovers to roughly the same stress level as before charging. The on-and-off cycle of the charging current can be repeated reversibly several times especially at high temperatures, e.g. 294 K, before fracture occurs. The fracture is exclusively intergranular as observed by scanning electron microscopy.
3.2.2. Below 190 K On application of the charging current the flow stress decreases only temporarily and then increases steeply (Fig. 1). The maximum stress is u n k n o w n since fracture occurs pre300,
250 RRRH ; 3 6 0 0 - 5 2 0 0 ~=
= 8.5 x IO-~/s ON ( 2 0 A / m 2)
r oF~-
200
~/~L 200
...-1 n
n
150
~150
_.m
ioo
~, IOO
50
50
,
,
,
,
I
,
i
,
I Tensile
,
I
i
2 Strain
,
,
,
I
i
i
:3 (%)
Fig. 1. T h e e f f e c t o f hydrogen charging on the flow stress o f h i g h p u r i t y i r o n at various l o w t e m p e r a t u r e s .
Tensile Stroin
Fig. 2. T h e e f f e c t o f hydrogen charging on the flow stress o f high purity iron deformed a t 170 K at three d i f f e r e n t s t r a i n rates.
210
maturely. Figure 2 shows the stress-strain curves at 170 K at three different strain rates but with the same charging current density. Fracture also occurs prematurely here but the fracture stress rises to 290 MPa at the two strain rates larger than 8.3 × 10 - 5 s -z . Hence we consider that at 180 and 170 K hydrogen charging causes hardening and that the softening immediately after application of the current is only transient. It is not known whether hardening would be observed at 190 K if fracture did not occur. Figure 3 summarizes the temperature effect for the grade A and B (RRRH = 3 5 0 0 - 5200) specimens. The upper curve represents the flow stress immediately before application of the charging current and the lower curve is the minimum flow stress during charging. The amount of stress decrease A o (the difference between the two curves) is largest near 200 K and decreases at higher temperatures. The solid symbols in the figure represent the fracture stress. The effect of hydrogen charging
changes drastically from softening to hardening as the test temperature decreases below about 190 K. It should be noted here that the temperature dependence of the flow stress above 200 K is markedly decreased by hydrogen charging.
3.3. The effect of purity Figure 4 shows the effect of hydrogen charging on the stress-strain curves of the grade C specimens (RRRH ~ 1800). At low temperatures the effect is similar to that in higher purity specimens, although the amount of stress decrease IA o l is smaller for the grade C specimens. At higher temperatures a small hump appears immediately after application of the current; the flow stress then decreases rather gradually. The hump grows as the temperature is raised and eventually at 273 K no decrease in the flow stress is observed. Thus near room temperature only hardening is observed. Softening still occurs in these specimens if a large current is applied, e.g.
300 = 8.3x RRR. o 0.1% v 250
300
10-5Is
~ 3500-5200 stress
flow
stress
just
before
charging minimum during
flow
fracture
stress
@
fracture
slress
1.7
OFF
charging
•
=
ON (20A/mz)
I 250
stress
(20A
200
I
/m z ) 170K o1
200 ~
x 10-4/s
o 150 CL
/
I
t
90K
o
150
v ~o O~ 03 O3
__~ I 0 0 ¢
I00
80A/m z
297K
50 50
RRR. E ,
,
,
,
200
,
,
,
i
i
,
,
250 Temperature
I
;
1500-1800
= 8.5 x 10-5Is
,
5(X) (K)
I
2 Tensile
Fig. 3. The amount of softening due to hydrogen charging at various temperatures. Hardening and fracture take place below 190 K.
Strain
:5
('/.)
Fig. 4. The effect of hydrogen charging on the flow stress of low purity iron with R R R H ~ 1 5 0 0 .
211
80 A m-2 as shown by the curve at the bott o m of the figure.
3.4. Electron microscopy Foil specimens (0.12 mm thick, 2 mm wide, 20 mm long) of grade B were tensile tested mainly at room temperature and at 200 K with and without hydrogen charging. TEM specimens were prepared by jet electropolishing from the foil tensile specimens and were examined in an electron microscope operating at 100 kV. Figure 5 shows a pair of electron micrographs of specimens deformed at 200 K. Figure 5(a) shows an uncharged specimen (e ~ 3.8%) and Fig. 5(b) a charged specimen (e ~ 3.0%, ic = 8 A m-2) *. The dislocation density is considerably increased in the hydrogen-charged specimen. By surveying several TEM specimens over a wide area it was found that in uncharged specimens the dislocation density varies from 3 × 10 s to 6 × 109 cm-2 from grain to grain whereas in charged specimens it varies in the range (3 - 6) X 109 cm -2 . The average dislocation density is larger by a factor of 5 - 10 in the charged specimens than in the uncharged specimens. The specimens deformed at room temperature with and without charging show no appreciable difference in dislocation density, which is in the range (1 - 5) × 108 cm -2 . Accurate determination of the dislocation character was impossible because none of the grains were found to be of a suitable orientation for a tilting experiment. No significant difference was found in the general features of the dislocation configuration between uncharged and charged specimens.
I10
g -- 2 0 0 I Fm
ITO
(a)
g = 200 4. D I S C U S S I O N
4.1. The mechanism o f softening The most remakable effect of hydrogen charging is the softening in the high purity specimens {grade A and B) at all temperatures above 190 K and in the less pure grade C specimens below room temperature. Several possible mechanisms for the softening have
* F o r a foil specimen the current density o f 8 A -2 m gives an a m o u n t o f softening similar t o t h a t in wire specimens. The effect o f specimen size on the s o f t e n i n g is r e p o r t e d in a separate paper [ 1 3 ] .
I Fm
(b) t
I
Fig. 5. T r a n s m i s s i o n e l e c t r o n m i c r o g r a p h s o f h i g h p u r i t y i r o n s p e c i m e n s d e f o r m e d at 2 0 0 K: ( a ) ' w i t h o u t h y d r o g e n charging, e ~ 3.8%; ( b ) w i t h h y d r o g e n charging, i c = 8 A m - 2 , e ~ 3.0%.
already been discussed to show that the softening is probably due to an interaction between hydrogen and screw dislocations [6]. However, the previous discussion was not necessarily complete, and a further and more
212 complete discussion is given below on the basis of the present new results.
4.1.1. Damage caused by hydrogen charging Softening may occur if the specimen crosssectional area is reduced b y the formation of cracks as a result of hydrogen gas precipitation. However, this mechanism is ruled out because the softening is reversible on switching the current on and off. Reversible softening by hydrogen charging is observed in mild steel and electrolytic iron when the charging current density is large and/or the specimen is thin enough [3, 1 2 ] . The softening is considered to be due to dislocation generation at damage, e.g. at blisters and internal cracks [2, 3, 1 2 ] . Hydrogen gas pressure built up in the damage during charging m a y enhance the dislocation generation. There is evidence to rule out damage as the cause of the softening in high purity iron. In the present research, blisters are formed in the grade C specimens which show hardening, not softening, at room temperature. The grade C specimens show softening at lower temperatures but the amount/x o of softening is smaller than in the grade B and A specimens in which very few or no blisters are formed. It is reasonable to assume that the number, if any, of internal cracks is also smaller in the grade A and B specimens than in the grade C specimens. Moreover, it might be difficult to explain the hardening observed below 190 K using the damage hypothesis.
4.1.2. Surface effects Some electrochemical reactions occurring during charging could cause softening, as in the case of the Rehbinder effect [ 1 4 ] . However, this mechanism is ruled out because very similar softening occurs in specimens {grade A) charged at room temperature and then quenched to and tested at low temperatures. In this experiment specimens were transferred from the electrolytic cell into a nonelectrolytic bath at low temperatures. Details of this series of experiments, which will be referred to as the "after-charging experiment", will be reported in a separate paper [ 1 5 ] . The softening could be caused by a steep solute concentration gradient. This t y p e of softening has been suggested by Sethi and Gibala [16] for the softening of surface oxide films. Their mechanism assumes that a large stress
produced by a steep concentration gradient at the surface nucleates edge dislocations, which are mobile under a much lower stress than screw dislocations. In the present experimental conditions the fugacity of hydrogen on the surface is of the order of 106 - 101° Pa [17] and the corresponding lattice solubility of hydrogen at room temperature is estimated to be 4 - 400 at.ppm by Gonzalez' measurement [5]. This solubility may give a hydrogen concentration gradient of 400 at.ppm per atomic distance at the most. In contrast, the hydrogen concentration gradient required to nucleate edge dislocations is estimated to be about 5 at.% per atomic distance [16] for a partial molar volume of hydrogen in iron of 2.0 cm ~ (mol H) -1 [18]. Thus this mechanism cannot be the cause of the present softening. Moreover, in the "after-charging experiment" mentioned above, the concentration gradient, if any, has probably disappeared before the tensile test is started because the fast diffusing hydrogen escapes from the surface layer immediately when the current is switched off. Nevertheless, softening is observed.
4.1.3. Bulk effects The above discussion together with those in previous papers [6, 7, 19] enables us to conclude that the softening is a bulk effect. In the following, several mechanisms will be discussed and it will be shown that the softening is most probably due to a hydrogenscrew dislocation interaction; hydrogen trapped at the core of screw dislocations increases their mobility.
4.1.3.1. The scavenging effect. Hydrogen might scavenge impurities already present in iron and remove the hardening effect of these impurities. This mechanism is readily discarded since, as stated in Section 3.3, the purer specimens show more pronounced softening. 4.1.3.2. Dislocation punching from precipitates. In a Va metal-hydrogen system hydride precipitates punch out edge dislocations, which are much more mobile than the screws, and cause softening [20]. It might be possible for hydride or hydride-like precipitates to exist in iron under an extraordinarily high effective pressure of hydrogen pro-
213 duced by charging. If such precipitates do exist, they may cause softening. Baranowski has suggested that iron hydride will exist, if at all, only under an extremely high hydrogen fugacity which can be estimated to be larger than 1021 Pa [21]. The effective hydrogen fugacity in the present experiments is considered to be less than 101° Pa. Hence the hydride hypothesis is unlikely. The result of the "after-charging experiment" is also unfavorable for the hydride hypothesis, as reported separately [15].
4.1.3.3. Increased dislocation density. Since TEM observation shows that the dislocation density in hydrogen-charged and deformed specimens is larger than in specimens without charging, we might consider that the softening is due to the increased dislocation density. As will be discussed below, the observed increase in dislocation density is not sufficient to explain the present softening. TEM observation shows that most of the dislocations are screws in both hydrogen-charged and uncharged specimens. The strain rate sensitivity of the flow stress at 200 K without charging has been found to be 20 - 25 MPa for a change in the strain rate by a factor of 10 [22]. The dislocation density is found to increase by a factor of 5 - 10 with hydrogen charging. This increase may cause a flow stress decrease of 10 - 25 MPa at 200 K if the strain rate is kept unchanged. The observed decrease in stress due to hydrogen charging at 200 K is about 80 MPa, which is much larger than the estimated value. To cause the observed softening the dislocation density would have to, increase by a factor of 103 - 104 . However, if the dislocation density were increased by this a m o u n t , it would cause an appreciable increase in the internal stress and hence hardening. Thus it is clear that the observed increase in mobile dislocation density is certainly insufficient to explain the observed softening. 4.1.3.4. The interaction of hydrogen with screw dislocations. The above discussion rules out all mechanisms for softening proposed so far except one, i.e. the interaction of hydrogen with screw dislocations. Solution softening and hardening have been observed in the iron-carbon [23] and iron-nitrogen systems [24] as well as in irradiated iron [25].
The softening and hardening have been discussed in terms of the interaction of screw dislocations with solute atoms and point defects [26]. The temperature range of the softening and hardening due to carbon and nitrogen is very similar to that for the present softening and hardening due to hydrogen. The similarity, however, does not mean that the same mechanism is operating for hydrogen as for carbon and nitrogen, since the behavior of hydrogen in iron in this temperature range is very different from that of carbon and nitrogen. Carbon and nitrogen are practically immobile below room temperature where they cause softening. Radiation softening is observed at temperatures where radiationinduced point defects are considered immobile. The hydrogen diffusivity, in contrast, is quite large (i.e. 6.3 × 1 0 - 6 c m 2 s - 1 [27] ) even at 200 K and hydrogen is considered to move together with dislocations. Hence theories of softening in which immobile solute atoms dispersed in the matrix enhance the double kink nucleation on screw dislocations by an elastic interaction cannot be applied to the present case. Since hydrogen is quite mobile it should be assumed to enter the dislocation core and m o d i f y the core structure to increase the kink nucleation rate. Hydrogen could be immobilized by forming hydrogen pairs, h y d r o g e n - i m p u r i t y pairs etc. which could act as stress centers to facilitate kink pair nucleation. This is unlikely, however. A resistivity study of the diffusion of hydrogen in high purity iron (RRRH >~ 3000) showed that excess hydrogen in the matrix escapes from the surface below 170 K and gave no evidence for the presence of a detectable number of immobile pairs in the matrix above 200 K [11]. Nevertheless, the "after-charging e x p e r i m e n t " reveals softening, as described in ref. 15. This observation rules out the hypothesis that immobilized hydrogen exists in the matrix and causes the softening. Moreover, if h y d r o g e n - i m p u r i t y pairs were the cause of softening, more softening should have been found in the less pure specimens; in fact the opposite was observed.
4.2. Proposed mechanisms for the softening and hardening If hydrogen atoms bound to the screw dislocation core are assumed to increase the double kink nucleation rate and hence the
214 screw dislocation mobility, the softening is readily explained. The hardening below 190 K is explained in terms of the hydrogen-edge dislocation interaction. Kinks on a screw dislocation move along it much more quickly than the screw dislocation itself. If the diffusivity of hydrogen atoms along a screw dislocation is sufficiently reduced below 190 K, hydrogen atoms may act as dragging points to the kinks. Since the Peierls potential of an edge dislocation is much smaller than that of a screw dislocation, hydrogen atoms may pin an edge dislocation rather than increase its mobility by reducing the Peierls stress. Thus below 190 K kinks are easily formed with the help of a hydrogen atmosphere on a screw dislocation but their sideways motion is hindered by hydrogen; hence hardening results. There is evidence that hydrogen pins the kinks (see ref. 15). The present experimental results, other than those discussed above, will be explained in terms of the proposed mechanisms. The transient softening observed at and below 180 K may be due to competitive operation of the softening and hardening mechanisms. For example, the following process may be suggested. The softening occurs at a relatively small concentration of hydrogen atmosphere because the presence of only a few sites of easy kink nucleation facilitates the motion of a screw dislocation as a whole. The hardening requires more hydrogen, since the easily nucleated kinks must be pinned before they move a considerable distance along the screw dislocations. According to this mechanism, softening will be observed first on application of the charging current, and hardening will take over when a sufficient amount of hydrogen has entered into the specimen. The result that more softening is observed at lower temperatures may partly be due to larger hydrogen concentrations in the atmosphere at lower temperatures. However, this effect would not be very appreciable. The atmosphere concentration is proportional to the concentration in the matrix, which is related by Sieverts' law to the effective pressure surrounding the specimen, at least at temperatures not t o o far below room temperature. Since the heat of solution of hydrogen in iron is negative, as the temperature is lowered, the matrix concentration is also
lowered. Consequently, the hydrogen concentration in the core will not increase appreciably as the temperature decreases. It is more likely that the temperature dependence of the softening is due to the reduction of the activation energy for the screw dislocation motion. In the above discussion it is assumed that the electrolytic reaction and the effective hydrogen pressure on the surface at a constant charging current density ic are independent of the temperature. This assumption is considered to be reasonable for two reasons: (a) similar A o versus i¢ relations are obtained at various temperatures; (b) a yield stress versus temperature relation similar to that in Fig. 3 is obtained in the "after-charging experiment". 4.3. The effect o f p u r i t y The concentration of dissolved impurity atoms is estimated to be about 10 at.ppm in the grade A and B specimens. The hydrogen concentration in the matrix during charging is estimated to be several tens of atomic parts per million [11] and it exceeds the impurity concentration. Hence the softening and hardening in these high purity specimens are considered to be the effect of hydrogen itself. The impurity concentration in the grade C specimens may be comparable with the hydrogen concentration, and it is reasonable to consider the hardening in these specimens at and above 273 K to be due to interactions of hydrogen and impurities with dislocations. As the temperature is lowered, softening, for which the mechanism is the same as for the high purity specimen, becomes predominant. The grade C specimens show softening at 297 K with increasing charging current density. This is probably due to the hydrogenenhanced screw dislocation motion, as in the grade A and B specimens. Since the concentration of hydrogen is larger for the larger current density, the hydrogen concentration charged with 80 A m -2 (Fig. 5) may be considered to exceed a critical concentration determined by impurities, and hence softening results. Blisters are formed in this charging condition (80 A m -2) but they will n o t cause appreciable softening. It should be mentioned here that a current density larger than 300 A m - 2 is necessary to cause damage-induced softening in an electrolytic iron with a similar electrolytic solution [ 1 2 ] . The degree of damage is much larger in an
215
impure iron than in a high purity iron. Hence damage caused by charging with 80 A m - 2 produces hardly any softening. 4.4. The fast response of the flow stress change The flow stress responds quickly to switching on of the charging current because of transportation of hydrogen by moving dislocations. Hydrogen also diffuses in towards the specimen interior by bulk diffusion. Although it is n o t known at present what fraction of hydrogen is transported by moving dislocations until the flow stress attains the steady level, the dislocation transportation should be dominant in the initial period of flow stress decrease. 4.5. The binding energy between a hydrogen atom and the core of a screw dislocation In the above discussion an appreciable binding energy is assumed to exist between hydrogen and the screw dislocation core. The binding energy between a hydrogen atom and an edge dislocation has been estimated to be a b o u t 0.3 eV [ 2 8 ] . No measurement has been performed, however, for screw dislocations. Chou has estimated on the basis of an elasticity theory that the binding energy of a solute atom with screw dislocations is about one-sixth of that with edge dislocations in iron [ 2 9 ] . Since the difference in the binding energy with the dislocation core will not be as great between edge and screw dislocations as the difference estimated from elasticity, the binding energy between a hydrogen atom and a screw dislocation core may be 0.1 - 0.3 eV. If hydrogen pairs are formed in the dislocation core (not in the matrix), the binding energy with dislocations will be increased. A binding energy of 0.2 eV is enough to form an appreciable concentration of hydrogen atmosphere at and around the screw dislocation core below room temperature. 5. CONCLUSIONS
Tensile tests were performed on high purity iron during hydrogen charging. The flow stress is increased by hydrogen charging below 190 K and decreased reversibly above 190 K. Hydrogen charging causes no, or only very few, blisters on the surface of specimens with RRRH > 3600, but an appreciable number of blisters on less pure specimens (RRRH
1800}. The less pure specimens were hardened rather than softened by hydrogen charging at room temperature. Several mechanisms for the hydrogeninduced softening, including damage due to charging, have been discussed, and it is concluded that the softening in high purity iron is due to an interaction between hydrogen and screw dislocations. The hardening below 190 K is due to an interaction between hydrogen and edge dislocations, including kinks on screw dislocations. The softening and hardening are the inherent effect of hydrogen on the mechanical properties of pure iron. It is proposed that hydrogen atoms at the core of screw dislocations m o d i f y their Peierls potential and thus increase their mobility. ACKNOWLEDGMENTS
This research was partly supported financially by the Ishihara-Asada Fund of the Iron and Steel Institute of Japan (S.M.) and by the Matsunaga Science Foundation (H.M.). The authors are grateful for this support. The authors thank Dr. S. Takaki for his cooperation in preparing the high purity iron specimens and Mr. Hosoya for his hydrogen analysis. REFERENCES 1 A. M. Adair, Trans. Metall. Soc. AIME, 236 (1966) 1613. S. Matsuyama, in Mechanism o f Delayed Failure o f Steels due to Hydrogen, Iron and Steel Institute of Japan, Tokyo, 1975, p. 113. Y. Tobe and W. R. Tyson, Scr. Metall., 11 (1977) 849. 2 S. Asano and R. Otsuka, Scr. Metall., 10 (1976) 1015. 3 S. Asano and R. Otsuka, Scr. MetaU., 12 (1978) 287. 4 C. D. Beacham, Metall. Trans., 3 (1972) 437. I. M. Bernstein, Scr. Metall., 8 (1974) 343. 5 O. D. Gonzalez, Trans. Metall. Soc. AIME, 245 (1969) 607. 6 H. Matsui, S. Moriya and H. Kimura, Proc. 4th Int. Conf. on Strength o f Metals and Alloys, Nancy, France, 1976, p. 291. 7 H. Kimura, H. Matsui and S. Moriya, Scr. Metall., 11 (1977) 473. 8 S. Takaki and H. Kimura, Scr. Metall., 10 (1976) 701. 9 S. Takaki and H. Kimura, Scr. Metall., 10 (1976) 1095.
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