Type-321 stainless steel is a Ti-stabilized steel which is designed to offer better ...... Bernstein and Anthony W. Thompson, Trans-AIME, (1981) P527-538. 32.
HYDROGEN CHARGING OF AISI-321 AUSTENITIC STAINLESS STEEL BY CATHODIC POLARIZATION
Oswald N. C. Uwakweh and Samuel Charca Department of General Engineering, College of Engineering P. O. Box 9044 University of Puerto Rico – Mayagüez Puerto Rico, 00681-9044
Vinod S. Agarwala Naval Air Systems Command, 48110 Shaw Road, Bldg. 2187/Suite 2373 Patuxent River, MD 20670-1906
ABSTRACT The room temperature cathodic polarization of cold rolled AISI-321 austenitic stainless steel leads to phase transformation as a function of charging medium. Of the investigated aqueous media, only the polarization with the mixture of sulfuric acid and arsenic oxide (0.1M H2SO4 + 1g/l As2O5 (IV)) resulted in the material degradation at current densities ranging from 500 A/m2 to 5000 A/m2. The extent of phase decomposition accompanying hydrogen charging depended more on charging time than on the prevailing charging current densities. Microscopically, formations of dislocations, stacking faults, culminating in martensitic transformation were observed based on combined X-ray diffraction and Electron Microscopic observations. The formation of macroscopic distortions and cracks along grain boundaries suggest that an inhomogeneous distribution of hydrogen was associated with steep hydrogen concentration gradients, and is used in rationalizing the observed features accompanying hydrogen ingress during the reported cathodic charging operations. Key words: Cathodic polarization, hydrogen charging, austenitic stainless steel, AISI-321, martensite, macroscopic deformation.
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INTRODUCTION The interactions of materials with hydrogen or hydrogen producing environments are well documented, and have been of considerable interests for a variety of reasons [1-5]. When metallic materials are exposed to direct hydrogen interaction, a phenomenon described as hydrogen embrittlement can occur at scales ranging from the atomic to macroscopic. The embrittling effects range from the alteration of mechanical properties as manifested in degradation or enhancements (tensile and fatigue strengths, etc) which can lead to catastrophic failures of high strength materials with serious consequences. Generally, steels are among the materials that are most disadvantageously affected by hydrogen because they are widely utilized. High-strength steels (HSLA, Aermet 100, 300M, and AF1410) and the stainless steels are extensively employed for aerospace applications in components such as landing gears and arrest shanks. In particular, those high-strength steels that have ferritic/martensitic microstructures [6-8] are widely affected by hydrogen producing environments. Stainless steels, especially of the austenitic type such as Type-300 series [9, 10], which are diversely employed in the chemical industries, are equally susceptible to hydrogen embrittlement. Type-321 stainless steel is a Ti-stabilized steel which is designed to offer better resistance to weld decay or heat affected zone is well associated with the formation of chromium carbides in welding applications. Because it offers more elevated temperature resistance, it finds use in aircraft exhaust manifolds and flanges and in jet engine parts. Their structural stability could be affected in the presence of hydrogen also. The problem of hydrogen interaction with materials leading to embrittlement is of high importance in the use of materials in diverse engineering applications. This explains why a number of different and varying approaches have been adopted in the investigation of hydrogen embrittlement, and their hydrogen transport behaviors. Diffusion and general transport properties of hydrogen in steels have been mostly based on the permeation technique originally developed by Devanathan and Starchurski [11], and later adapted by De Luccia et al. [12, 13]. These methods have been applied in the studies of different classes of steels such as the high-strength stainless and duplex stainless steels. In addition, there is a simple electrochemical charging method in which atomic hydrogen is produced from electrolyses of appropriate aqueous media to study steel/-hydrogen interaction and hydrogen embrittlement. Based on electrochemical processes involving cathodic polarization, a range of steel materials have been studied based on their abilities to absorb hydrogen during cathodic polarizations in aqueous media. As a consequence, the resulting hydrogen charging of these steels followed hydrogen ingress that lead to alloy degradation such as in the cases of austenitic stainless steels [14-17] and high carbon Fe-C austenitic alloys [18] are well documented. In order to understand the mechanisms by which the decomposition of austenitic steels is accomplished, varied interpretations are documented [14, 16, 1923]. Similarly, while there are different ways of introducing hydrogen in a steel material. The room temperature charging of austenitic alloys continues to be a valid simple means of investigating how their properties are affected by hydrogen and phase stabilities with the aim of deriving or arriving at a coherent explanation towards their individual hydrogen embrittlement behaviors. In this study, we report (to the best of our knowledge) on the first account of the macroscopic deformations accompanying room temperature charging of as-received cold worked Type-321 austenitic stainless steel in an electrolytic bath consisting of dilute acid or basic solutions impregnated with different hydrogen recombination inhibitors otherwise known as poisons or atomic hydrogen promoters. EXPERIMENTAL Samples were cut from cold rolled austenitic steel foils of Type 321, which was supplied by SHOPAID Inc., Woburn, MA 01801. The samples were cut from the 0.002” (0.5mm) cold rolled foils as
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rectangular sheets with dimension in the order of 2 cm x 10 cm followed by initial degreasing, and cleansing using ethanol, and subsequent air drying. The samples were used in their as-received surface finish characterized by a dull shine. The samples were cathodically charged in aqueous electrolytic media impregnated with hydrogen promoters, or “hydrogen recombination reaction poisons” necessary to aid hydrogen ingress in the austenitic steel. The composition of Type 321-austenitic steel is given in Table 1 below. TABLE 1 COMPOSITION OF AISI-321 AUSTENITIC STAINLESS STEEL Chemical Composition (Wt Pct) *=max / Bal = Fe AISI-321
C* 0.08
Mn * 2.00
Si * 1.00
Cr 17.00/19.00
Ni 9.00/12.00
Timax = 0.7 5 x %(C+N)
S* 0.03
P* 0.045
The electrolytic media investigated in this study are given in Table 2 with their corresponding concentrations and poisons. All the charging operations were carried out at ambient temperatures under a fume hood. This implied that the evaporation rate of the electrolyte was constantly checked in order to maintain same level of electrolyte through out the charging period. The samples were held suspended in the electrolyte and platinum counter electrodes were used as an anode to complete the electrical circuit. All electrochemical charging were carried out under constant current (galvanostatic) modes. This enabled charging current densities to vary from 160 A/m2 to 200 A/m2 for times varying from 3 hours to up to 96 hours. It was important that the contacts between the steels and the DC power supply were not affected by the charging media, i.e., portion of the steel connected to the DC power supply was not exposed to the electrolytic charging medium.
Electrolyte
TABLE 2 COMPOSITIONS OF THE ELECTROLYTES INVESTIGATED Concentration Poison type Electrolyte Overall Composition
NaOH H2SO4 H2SO4 H2SO4
0.1 M 1N 0.1 M 0.1 M
NaCN (NH2)2CS As2O5(V) Na2HAsO4.7H2O
0.1M NaOH + 0.2M NaCN 1N H2SO4 + 1g/l (NH2)2CS 0.1M H2SO4 + 1g/l As2O5 (V) 0.1M H2SO4 + 1g/l Na2HAsO4.7H2O
After hydrogen charging, the samples were rinsed in water followed by a methanol rinse, and then airblow dried. The samples were subsequently subjected to X-ray diffraction studies at room temperatures. These measurements were carried out using the Rigaku-5000 Diffractometer, with Cu Kα radiation at 40 kV and 35mA. Since the diffractometer did not have low temperature measurement capability, no attempts were made at sub-ambient measurements. The scanning range of the x-ray diffraction measurements was chosen to capture all possible diffraction peaks, and so, in the range of 20° ≥ 2θ ≤ 120°. Next, thin foils of hydrogen charged stainless steel specimens were prepared for Transmission Electron Microscope (TEM) using a Gatan Dual Ion Mill at a gun current and gun voltage of 1 mA and 7 kV respectively. A JOEL JEM-100CX II TEM operating at 120 kV was used for the observation and analyses of microstructures of the specimens both prior to and after cathodic charging at ambient temperatures. Given that more than 24 hours was generally required to obtain a hole, overall stability of the charged alloy was more important in this study than the focus on investigating time dependent structural variations as was conducted by Shuchen et al [19]. Subsequent examinations of the foils were conducted with the use of a JOEL JEM-100CX II microscope (TEM) operating at 120 kV.
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Hydrogen content was not determined in this study, since the focus of the study was to determine the long-term effect of cathodic charging with the aim of determining its relative stability with respect to charging media degradation effects. Future studies will focus on reconciling the amounts of hydrogen ingress with respect to permeation properties, and subsequently its overall degradation mechanisms.
RESULTS AND DISCUSSION Effects of cathodic polarization medium on: Hydrogen charging, Phase decomposition, and Macroscopic Deformation of AISI-321 austenitic stainless steel When the present samples were hydrogen charged via cathodic polarization in a 0.1M NaOH + 0.2M NaCN solution they did not exhibit any macroscopic distortion or changes. Similarly, no macroscopic change was observed with the charging of the Type-321 austenitic stainless steel in a 1N H2SO4 + 1g/l (NH2)2CS solution. However, in the second case a black thin layer was observed on the surface of the steel. The situation was much different when the solution consisted of 0.1M H2SO4 + 1g/l As2O5(V) under similar charging current densities. The use of different charging electrolytic media was based on the fact that they have been used in the study of other steels. For example, 0.1M NaOH + 0.2M NaCN solution was used in the study of hydrogen permeation properties of Armco- Iron and the ultra high strength AF1410 steel [12, 13]. On the other hand, 1N H2SO4 + 1g/l (NH2)2CS was used in the study of hydrogen diffusivity in Aermet 100 high-strength steel using galvanostatic charging [24] techniques. In various characteristic permeation studies of duplex stainless steels, the charging media was reported to be simply 0.1M NaOH [25-27] with the hydrogen charging lasting in some cases for up to two weeks (28 days, i.e., 4 weeks). In the charging of Types 304 and 305 stainless steels for varying times, use of 1N H2SO4 + 0.25g/l NaAsO2 was reported by Chen et al. [19], while Young et al. [28] employed 10 ml H2SO4 + 1000 ml H2O + 0.8g sodium pyrophosphate in the study of precipitation hardened stainless steel. Lastly, mild steels and maraging steels of grade 350 were reportedly charged with 1wt% NaOH + 1g/l As2O3 by Tiwari et al. [29]. Since the samples in the Tiwari et al. [29] study were bulk tensile test pieces under pre-stress conditions, it was not expected that any macroscopic distortions would arise. Given the disparity in the steels and employed charging media, we deemed it necessary to initiate a systematic study in an attempt to determine the effects of charging media on the stability and hydrogen transport behavior of steels using AISI-321 as our initial test case. The connection between the macroscopic distortion of austenitic materials subjected to hydrogen charging via cathodic polarizations are rather limited [18, 30], while their decomposition to other product or new phases abound. For instance, in the charging of Types 304 and 305 stainless steels for varying times in 1N H2SO4 + 0.25g/l NaAsO2 Chen et al. [19] reported the occurrence of phase transformations associated with the sequential formation of hexagonal ε, εH, (hydrogen dissolved in the ε phase) and α’- martensitic phases for the type 305, while a hydride phase designated as γ* was reported to have formed with Type-304 prior to the martensitic phase. There was no unique mention of the buckling of the sample used by these authors, most probably on account of the relatively small sizes of their experimental samples which were 7-mm2 x 1-mm thick, even though they reported charging times of up to 100 hours. On the other hand, Young et al. [28] employed 10 ml H2SO4 + 1000 ml H2O + 0.8g sodium pyrophosphate in the study of precipitation hardened stainless steel, and reported neither macroscopic deformations nor distortion of the charged materials. In contrast, Hänninen et al. [31] reported the buckling of thin sections of transmission electron microscopic (TEM) samples after hydrogen charging for types-310 and 316 stainless steels.
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The cathodic charging of Type-321 austenitic stainless steel leads to macroscopic deformation when carried out continuously for more that 24 hours. While the early stage of the deformation was entirely microscopic, the long-term effects of charging can be observed visually as the sample distorts with charging time. Figure 1a, shows the photograph of a typical sample charged for up to 96 hours using a 0.1M H2SO4 + 1g/l As2O5(V) solution. It can be observed that the damage, as evidenced by the bending of the sample, is extensive and becomes very severe where tearing (or rumpling) of the sample can be observed as shown in Figure 1b. It was observed that while macroscopic deformation occurred consistently following extended charging times, the orientation of the deformed surface with respect to the exposed un-immersed portion varied from one charging operation to another. This is not surprising because the uncharged starting material being polycrystalline implies that there will be some level of statistical variations to be observed from one sample to another, while the overall effect would remain unchanged. At the microscopic level, it is expected that the structural changes would be similar from one charging operation to the other, as will be shown in subsequent measurements. Further, with the observation of same trend in the macroscopic deformation of the AISI-321 steel in the electrolytes containing arsenic- based poisons (due to their effectiveness in facilitating hydrogen ingress), we proceeded to test masked samples. This was carried out in order to determine or elucidate the observed behavior as a function of structural change and hydrogen ingress. The masking was carried out with a Lacquer, such that two sets of samples had exposed surface on one side, and a limited exposed area surface on the rectangular sheets were realized. Both types of masked samples exhibited similar behavior with respect to their macroscopic deformation. As shown in Figures 2a, 2b, 2c, and 2d, these are macroscopic deformations were indeed as a result of hydrogen ingress and not due to some mere surface phenomena. The room temperature aging of these masked samples also show that hydrogen escape or redistribution or portioning after cathodic polarization led to the partial restoration of the macroscopic deformation observed. Figures 2a and 2b showing three stages associated with the deformation of cathodically polarized samples illustrate that hydrogen diffusion is faster or enhanced in strained or deformed materials based on the observed profiles of the deformed material. One can notice bending over the top portion of the initially bent sample, while the lower portion depicts continued bending. These figures show that bending led to strained or stressed sections of the material, which in turn promoted accelerated hydrogen ingress and explains the observed tearing along the edge of the samples as a stress relieving process or measure. It was observed that while macroscopic deformation occurred consistently at extended charging times, the orientation of the deformed surface varied from one charging operation to another as shown in Figures 2d and 2e respectively. This is not surprising because the starting material was polycrystalline, and therefore statistical variation of surface defects at the microscopic level could affect initial local hydrogen distribution from one sample to another, with same overall effects conserved. At the microscopic level, it was expected that the structural changes would be similar from one charging operation to the other, as will be shown in subsequent measurements. Our findings show that care must be taken in choosing the appropriate medium for the purpose of hydrogen charging of steels by cathodic polarization. In this regard, we would like to draw attention to a prior study of hydrogen diffusivity in AISI-304 where the Devanatha and Starchurski double cell (DSDC) [11] arrangement was used by Owczarek and Zakroczymski [26, 27], and hydrogen charging was carried out for 14 days using NaOH solution. It is surprising that in our study, no macroscopic deformation was detected after such a long time of charging. On the other hand, our experience with the cathodic polarization of austenitic materials in general show that even with very small polarization currents, hydrogen ingress would eventually lead to macroscopic deformations. According to Uwakweh and Genin [18], charging of Fe-1.95C resulted in the macroscopic distortions of the austenitic samples to a high degree comparable to the case of type 321-stainless steel of this study. It was noticed that
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martensitic transformation occurred with buckling, while the extent of the buckling correlated with the volume of martensite formed. Given the high carbon content of the binary alloy, the resulting martensitic phase being tetragonal was designated as α’T to distinguish it from the body centered cubic phase martensite (α’) that results from the decomposition of the stainless steel. X-Ray Diffraction Measurements The room temperature X-ray diffraction measurements conducted on the as-received and hydrogen charged samples over the 20° ≤ 2θ ≤ 120° range are shown in Figures 3a and 3b respectively. The spectral peaks as detected in the measurements are tabulated in Table 3. TABLE 3 X-RAY DIFFRACTION (XRD) PEAKS POSITIONS Peaks positions in the As-received /Charged samples in 0.1M H2SO4 + 1g/l As2O5 (V) Sample As-Received Charged
Centroid Spectral Peaks Positions (2θ°) 37.0 – 44.0 47.0-51.5 55.5-75.0 75.0-90.0 47.909 (200)γ 37.844, 41.461, 47.535, 43.629 47.617 43.698 40.715 (111)γ
71.785 (220)γ 58.643, 58.699 58.747, 61.788 71.628
87.788 (311)γ 79.678, 87.865
92.5100.0
95.895
Correspondingly, Figures 3a and 3b respectively show the spectra aligned in the same scale for easy comparison. Being an austenitic face centered cubic structure, the spectral peaks were indexed in the following order (based on the extinction principles): (111), (200), (220), (311), (222), (331), (420), (422), (333)/(511)). It can be seen on Figure 2a that only the first four peaks were observed while the others were missing. In the as-received state, the peaks were seen to be fairly sharp with the entire spectrum showing only austenitic peaks. Table 3 shows the different peak positions in terms of their angular positions. Next, there were marked differences in comparing the spectra of the as-received and charged states of the steel, especially in the low angle regions. Apart from the new peaks, the original peaks associated with the as-received material can be seen to assume different profiles including broadening, and slight positional changes. These observations are important in the context of the fact that the measurements were carried out on ambient temperature aged materials following cathodic polarization. Given that hydrogen concentration gradients develop with charging, it is to be expected that there will be some alteration of the microstructure of the charged material at the subsurfaces. The peak positions identified are grouped in Table 3. In the aged state, it is expected that hydrogen concentration gradient build-up which normally accompanies cathodic polarization would have eased off, while the microstructural deformations accompanying this hydrogen concentration build-up would reach an equilibrium condition, with possible slight structural restoration.
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Peak broadening can be associated with strained lattice, development of microstructural defects such as dislocations, stacking faults, and voids, etc. However, due to the non-uniform concentration build-up and the antecedent micro-stress/micros-strain development within a few angstroms (Å) below the cathodically polarized surface where the hydrogen ingress took place, a redistribution of structural defects will result in explaining the spectral peaks broadening. Thus without the formation of stable hydrides, the non-uniform distribution of structural defects can both account for the observed peaks shifts, broadening, and changes in intensities. These peaks were indexed similarly to those reported in the literature, including recent neutron diffraction measurements on types 304 and 310 austenitic steels [19, 32]. The intensities of these peaks are affected by the rolling operation (during processing) which imparts some texturing to the sheet metal. In comparison to the as-received material, the phase decomposition accompanying cathodic charging is evident with the observed peaks in the ranges given on Table III. The shifts of the original austenitic peaks after hydrogen charging can be observed in the ranges listed on Table III. In addition, the emergence of new peaks as a direct result of the charging process shows that phase decomposition did take place which was coupled with structural changes as well. In the asreceived state, only the peaks identified on Table III were observed, and in addition to spectral peaks displacements, and emergence of newer peaks, one can also observe changes in peaks intensities compared to the as-received spectrum. Chen et al. [19] reported that the charging of metastable 304 and 305 austenitic stainless alloys (laboratory heats) resulted in the formation of hydride phases. While they showed only the spectral peaks in limited regions, rigorous comparisons can not be made with our work at this stage, since all our x-ray diffraction measurements were conducted after room temperature aging following cathodic charging. In addition, they concluded that while the transient hydrides transformed to ε-martensites, not all the ε-martensites transformed to stable α’-martensite at ambient temperatures. Thus, one can say that within similar regions, there was evidence of both the ε-martensite which we reconcile as stacking faults based on [19], Troiano and Whiteman [20], and Burke et al. [21] and the formation of more stable α’ martensite.
Electron Microscopic Examination Scanning Electron Microscopy (SEM) Figures 4a and 4b show typical SEM images of the charged surface revealing mud crack patterns. The pattern suggests presence of intergranular cracks more than the intragranular types. They show also that room temperature aging leads to hydrogen escape resulting in additional cracks development. Transmission Electron Microscopic (TEM) Studies Transmission electron microscopic (TEM) examination of the thin foils of the charged samples as displayed in Figure 5a shows the presence of thin martensite lathes with faults. Figures 5b and 5c show dislocated regions. The arrowed parts show dislocation loops such as the Frank-Reed types, which suggests multiplication of pinned dislocations, while Figure 5c shows region heavy dislocation network. The stacking faults can be seen in some cases emanating from grain boundaries, as shown in Figure 5e where they run in parallel. These images show that the structural changes accompanying the phase decomposition of Type-321 stainless steel were not limited to surface regions alone. This is attested by the presence of dislocations and stacking faults respectively. The presence of heavy dislocations networks suggests that the hydrogen charging leads to the formation of local strains, while the exact mechanisms associated with their development cannot be directly inferred with the manner of investigations conducted. In Figure 5b, we see in addition to a well defined dislocation, stacking faults and weak contrasts of needle-like α’-martensite which corroborates the x-ray diffraction results. This is because stacking faults can lead to satellite or slit spectral peaks associated with high atomic density 7
planes. The weak contrasts can also be attributed to non-flat TEM surface arising from the macroscopic distortion accompanying hydrogen charging of the Type-321 stainless steel. In Figure 5b, dislocations, stacking faults, and martensitic lathes are seen within the same austenitic grain. Close examination of the dislocations show features typical of the occurrence or presence of Frank-Reed type dislocation source as said earlier. Interestingly the advancing front of the dislocation loop can be inferred between the two dislocation loops. The change in orientation of the fault line suggests probable cross-slip which may have been associated with hydrogen transport. In relation to the macroscopic deformation shown above, one can explain the features of the TEM micrographs. For one, the buckling of the samples during cathodic charging was gradual, and eventually culminated to tearing through serration of the edges of the charged samples. Since the buckling observed was maximum at the mid section of the immersed portion (under cathodic polarization) it is to be expected that the relative difference between the hydrogen transport from the immersed and un-submerged portion created nonuniform local stress/strain which eventually led to microstrain, and subsequently to macroscopic strain, and tearing along the sample edges. It is important to note that stress-strain relief through hydrogen egress did not lead to surface pitting or tearing on all the charged samples. The suggestion that the formation of deformation-induced martensite is an important factor which controls hydrogen susceptibility of austenitic steels [33, 40] seems to be contradicted by the observed hydrogen cracking of stable austenitic phases [34-38]. Because most austenitic steels have low stacking fault energies, as is equally true for most FCC materials, it has been suggested that hydrogen embrittlement can be understood from the perspective of reduced stacking fault energy and the accompanying prevalence of planar slip increase [34, 36]. This view hinges on the idea that planar slip concentrates the hydrogen along slip planes, while cross-slip serves to dilute the hydrogen throughout the matrix material as hydrogen is taken to be dislocation bond while moving with the dislocation during slip. On the bases of hydrogen charging at ambient temperatures, the following structural decomposition sequences have been reported for AISI-304, 310, and 316 austenitic steels [38] thus: AISI-304 and 316: γ→ ε →α (with 316 reported to be more stable than 304 alloy) AISI-310 γ → γ* (expanded austenite) / ε* (expanded ε-martensite) In addition, the transformation γ → γ* (expanded austenite) is reported to be reversible at room temperature, while Kamachi [39] additionally reported the presence of metastable FCC-hydride and an intermediate hexagonal ε-phase. Based on extensive studies conducted on AISI-304, 310, and 316 steels, Hänninen et al. [31, 34] concluded that hydrogen induced embrittlement to be reversible. However, they conducted their study on bulk samples that were charged in aqueous acidic (1N H2SO4 + 0.25g/l NaAsO2). Their TEM observations were based on cut samples from either the outer surface or the internal portions of their tensile (0.3mm thick) samples with the microstructural features strongly dependent on location.
DISCUSSION As is evident, the entry or ingress of hydrogen in metallic materials poses a very serious problem given that degradation of materials properties occur. However, the problem is not limited to only hydrogen entry, but also in the mechanism of its entry. To this end, the Devenathan and Starchurski cell (DSDC) [11] has played significant roles in elucidating the synergistic relationship between hydrogen evolution reaction (HER), permeation flux, and the charging and hydrogen evolution (recombination fluxes) [21].
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As reported by Hoelzel et al. [22] the high pressure gas phase charging of Type-304 and -310 austenitic stainless steels generated large amounts of dissolved hydrogen without a corresponding macroscopic distortion. In the present study, hydrogen charging by cathodic polarization of the type-321 stainless steel implied the existence of different stages leading to or associated with the steel material hydrogen uptake. Firstly, atomic hydrogen must be liberated at the surface of the stainless steel, then some recombination of the atomic species to molecular hydrogen reflected by the presence of escaping gas bubbles, while some atomic specie diffused or dissolved in the metal material. These stages should not be kinetically similar in with respect to the gas phase and electrochemical methods. It is known that poisons such as the arsenic based ones aid hydrogen ingress during cathodic charging of austenitic steels and austenitic super-alloys such as the IN-903 (inconel 903) [41]. According to Robinson et al. [41], charging austenitic IN-903 at room temperature in unpoisoned and poisoned electrolytes resulted in dramatic difference in deuterium charging of the alloy. With the unpoisoned electrolyte, the surface concentration of deuterium increased from 3x103 appm (atoms parts per million) to 2x104 appm for charging current densities of 0.1 μA/cm2 to 10μ A/cm2 2respectively. Meanwhile, charging with poisoned (250 mg/l of NaAsO3) electrolyte resulted in surface concentration of 0.6 atom fraction at 10 μA/cm2 charging current density for comparable times. Though they did not report macroscopic deformation as in our case here, they never the less observed the correlation of development of slip bands, and matrix carbide cracking. The electrochemical method of charging therefore can be varied by controlling the medium to the extent that is not comparable with the gas phase charging, where temperature and pressure are the only process variables. In the case of electrochemical charging, one has a lot of latitude in the manner of hydrogen charging which include control of polarization current density, concentration and type of poison, concentration and type of acidic or basic medium for instance. In the gas phase case, since no evidence of molecular trapping is reported, the molecular gas must dissociate before eventual diffusion or dissolution into the metal material. Further, the dissociation, adhesion, and diffusion/dissolution stages or processes must have proceeded in such a manner that macroscopic distortions did not occur. The elastic and inelastic neutron scattering experiments did not report any formation of hydrides during the charging process, but only increasing Me-H (i.e., metalhydrogen) interatomic distances [29]. In addition, these authors did not detect the formation of εmartensites, and therefore concluded that the formation of ε-martensite might be determined by the presence of absolute hydrogen concentration rather than the stress states resulting from a particular hydrogen distribution, which contrasted the cases of hydrogen charging by electrolytic means. The gas phase charging is inherently different from the electrochemical one because one obtains uniform distribution of hydrogen without any known apparent rate limiting steps prior to hydrogen ingress as with the case of electrochemical charging. Defects introduction accompanying hydrogen is more with electrochemical charging than with the gas phase charging. From the observations made in this study, it is conceivable to imagine that apart from the stages associated with hydrogen uptake listed above, there could be an additional effect, such as the weakening of the atomic bonds in the steel material due to the presence of hydrogen recombination poisons at any stage in the charging process. This is because, only one electrolyte (0.1M H2SO4 + 1g/l As2O5(V) ) gave rise to macroscopic distortions of the stainless steel during charging. Therefore, it is possible that the combined action of weakening bond strength in the metal with hydrogen liberation/adhesion/and diffusion led to build up of hydrogen concentration gradients from the exposed surface to the subsurface and ultimately large strains leading to macroscopic distortion. The development of stress as a result of non-equilibrium diffusion of atomic species into a material is said to be tensile on egress while compressive on ingress according to Li [42]. Robinson et al. [41] used the result of his analysis to 9
determine that the threshold for cracking of IN-903 assuming a partial molar volume of hydrogen to be 1.73cm3/mole [43] was deuterium concentration in the range of 1.5x 04 and 1x105 appm. On the basis of this, we see that our hydrogen charging must have attained comparable levels of surface concentrations in order to provoke the cracks observed during room temperature aging of the charged AISI-321 samples as shown in the SEM micrographs of Figure 2. The x-ray diffraction spectral peaks of materials charged at varying degrees did not reflect the possibility of stable hydride phases being present during the x-ray diffraction measurements. Further, visual observation of the samples during charging showed that buckling of the cathodically polarized stainless steel material was gradual and continuous. If these were associated with hydride phase/compounds, such compounds capable of causing macroscopic distortion of the steel would have been detected after the charging processes. Thus, it is suggested that while the recombination poisons aid the overall hydrogen charging process, the uneven distributions of the ingressed hydrogen coupled with the weakening of the stainless steel material bonds would be the likely explanation for the observed macroscopic strain. There exists a correlation between type of stainless, nature of prevalent environment, and accompanying phase decomposition. For instance, studies indicate that Type-310 stainless steel does not exhibit phase changes when charged to achieve hydrogen contents ranging from hydrogen/metal ratio (i.e., H/Me) of 0.003 to 1.03 with gas phase charging, corresponding to the pressure range of 3.0 to 7.0 GPa. On the other hand, for Type-304 stainless steel, formation of ε-martensite was observed for hydrogen concentrations of H/Me=0.56 at 3.0 GPa and H/Me=1.03 at 7.0 GPa respectively. The formation of εmartensite was also determined in Type 304 stainless steel that was subjected to 4.0 GPa in the absence of hydrogen [22]. It is therefore important to recognize the marked difference exhibited by these stainless steels in the presence of different hydrogen environments. For instance, while it is known that martensitic transformation can lead to macroscopic deformations in steels, the levels of hydrogen ingress reported in the gas phase charging of Types-310 and -304 by Hoelzel et al. [22, 32] were not enough to provoke macroscopic distortions. This suggests that the charging media and method of charging influence the hydrogen injection process. Non-uniform distribution of ingressed hydrogen as opposed to uniform hydrogen distribution would likely cause macroscopic deformation as observed in our study. Next, given that hydrogen introduction into steel via electrochemical charging is introduces a lot of defects as evidenced from XRD peaks shifting and broadening, coupled with SEM and TEM micrographs, the propensity to form ε-martensite (which is associated with stacking faults) would form more readily than in the gas e charging. The study which reported distortion behavior of Types-304 and -310 austenitic steels during high temperature gas phase hydrogen charging [32] showed that εmartensite formation characterized only Type-304 after very high pressures, while none was observed with the Type-310 steel. Thus, independent of the stainless steel type, macroscopic distortions to the level or extent observed in this study have not been reported to the knowledge of the authors. In addition to the observations noted about the formation of ε- and α’-martensite in Type-304 and 310 stainless steels, Okada et al. [44] noted the development of surface cracking and the formation of εmartensite in Types-304 and 301 austenitic stainless steels. In their case, the ε-martensite formed in Type-304 steel persisted with room temperature aging, in contrast to the case of Type-301 steel where it decomposed with room temperature aging. In contrast to their observation, Narita et al. [15], identified only α’-martensite formation in their study of austenitic stainless steels. These conflicting accounts imply that there is room for further studies on these alloys. One can not however ignore the possible effects of alloy processing histories (as some materials used in these studies were laboratory heats, instead of commercially procured materials) on the observed behaviors.
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Concerning hydrogen transport, Farrel and Lewis [45] determined the hydrogen diffusivity of hydrogen in Type-310 stainless to be 1.4x10-12 cm2/s at 298K which they used to estimate surface concentration in the range of 0.5 to 0.8 hydrogen atoms/metal, while based on permeation experiments using DSDC [11] arrangement on Type-304 steel, Zackroczymski [46] determined a value of 7.5x10-5 cm2/s with a NaOH charging electrolyte. According to Zackroczymski, the diffusivity was close to the lattice diffusivity of hydrogen. The marked difference between the values reported by Farrel and Lewis [45], and Zackroczymski underscores the need for caution in making estimates based on reported data on one hand, and the need to investigate imparts of charging media on the hydrogen transport behavior in steels. Lastly, martensitic transformation is usually accompanied by stress/strain within the harboring matrix phase, and can lead to surface distortions as well. On the other hand, formation of stacking faults is not known to lead to the type of surface distortion or deformation typical of martensitic transformations. Based, on this, our observation of α’-martensite should be a contributing factor in the formation of macroscopic deformation observed in our charged samples. Next, the martensitic and austenite interface would become attractive site for hydrogen trapping, which could exacerbate hydrogen ingress. Coupled to the steep hydrogen concentration associated with cathodic charging, cracking observed with extensively or prolonged charging can be explained.
CONCLUSIONS The following conclusions can be drawn from the results of the conducted studies: 1. Phase decomposition accompanying cathodic charging of Type 321-stainless steel in aqueous media is a function of the charging media. Of the media investigated in this study, only 0.1M H2SO4 + 1g/l As2O5(V) was effective in leading to phase decomposition of the cathodically polarized austenitic phase, as opposed to the two namely, 0.1M NaOH + 0.2M NaCN, and (NH2)2CS solutions, respectively. 2. Phase decomposition leading to the formation of ε-martensite phase occurs in addition to α’martensitic phase. 3. Macroscopic distortion is observed during the charging of 0.002” (0.05mm) thick cold rolled Type 321-austeniticsteel foils. The macroscopic distortion can be associated with the development of weakening of steel atom-atom bonds, large and unequal concentration gradient distribution in the charged material, and eventual development of structural imperfections such as dislocations and stacking faults. 4. X-ray diffraction (XRD) measurements show (lack of) no evidence of stable hydride phase formation and lattice expansion of in the as-received material. Similarly, the formation of α’martensite in addition to ε-martensite or stacking fault can be deduced. 5. There seem to be a correlation between the nature of stainless and its interaction with hydrogen with respect to distortions.
ACKNOWLEDGEMENTS The authors would like to thank Dr. Rabina K. Mahapatra for his help in TEM, Mr. Kowalik, and Charles Lei for helping in the X-ray diffraction measurements, and William Lightell for taking photos of macroscopic deformations. In addition, one of the authors, ONCU wishes to acknowledge the support and guidance of Dr. Yapa Rajapakse, the program manager of ONR-grant # N000140310540, and support from US-Navy / ASEE Summer Faculty Fellow Research Program.
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(a) (b) Figure 1a: Intermediate stage of Macroscopic Deformation of AISI-321 during hydrogen charging. (b): Advanced stage of macroscopic deformation of AISI-321 during hydrogen charging.
(a)
(a)
(c)
(d)
(e)
Figure 2a: Uncharged masked AISI-321 sample. (2b): Curved portion of masked sample. (2c): Delaminated masked sample with hydrogen charging. (2d): Uncharged partially masked AISI-321 sample. (2e): Curved portion of partially masked sample with hydrogen charging
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Figure 3. (a): X-ray diffraction pattern of As-received AISI-321 and (b): Xray diffraction pattern of H-charged via cathodic polarization.
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(a) (b) Figure 4 (a): SEM micrograph of AISI-321after hydrogen charging (b): SEM micrograph of AISI-321 after hydrogen charging with room temperature aging
(a)
(b)
(c)
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Figure 5 (a): Martensitic lathes (arrowed), Stacking Fault region and high dislocations density areas of hydrogen charged AISI-321 (b): Martensitic lathes (unfilled arrow heads), Stacking Faults, and Dislocation loops (filled arrow heads) (c): Highly dislocated region (d). Stacking faults (parallel) within a grain around triple junction
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