5
Embrittlement of Amorphous Fe40Ni38Mo4B18Alloy by Electrolytic Hydrogen J.-J. LIN and T.-P. PERNG The effects of hydrogen on the tensile properties and fracture processes at room temperature were investigated. Specimens were tested at various strain rates in air or under different cathodic charging-current densities. The slopes of the stress-strain curves were essentially identical for all the specimens, except that the fracture points varied under different test conditions. Macro- scopically, hydrogen only affected the elastic deformation behavior, but microscopically, the embrittlement was caused by the heterogeneous nucleation of localized plastic deformation. The degree of hydrogen embrittlement increased as the charging current increased or as the strain rate decreased. With the same charging current and time, longer dynamic charging resulted in more severe embrittlement. Before fracture took place, the strength of the alloy could be com- pletely restored if hydrogen had been removed. Hydrogen diffusivity and solubility were used to draw the time-dependent hydrogen concentration profiles for the specimens under different charging conditions. The difference in the mechanical properties was correlated with the hy- drogen concentration within the specimen. I. INTRODUCTION IT has long been recognized that hydrogen may embrittle many crystalline metals, including Fe- and Ni-base alloys. The embrittlement is manifested by the nonductile fracture mode, reduced ductility, and reduced tensile strength in a tensile test. Recently, some Fe- and Ni-base amorphous alloys also have been found to ex- hibit similar hydrogen embrittlement (HE) in a tensile or bending test by cathodic charging in acidic solutions, f~-81 Amorphous alloys are distinguished from the crystal- line alloys, particularly in their deformation behav- ior. M9-~31 Macroscopically, they behave in a brittle manner. In a tensile test, the linear relationship of stress strain extends to the fracture point. Plastic deformation occurs only shortly before or simultaneously with frac- ture. Plastic strain is localized in sharp shear bands, c~~ Their fracture is preceded by large local plastic shear, which produces a featureless smooth zone, followed by a catastrophic shear, which produces a "vein" or "ridge" pattern. 1~~ The lack of ductility has been ascribed to the inhomogeneous deformation. Therefore, the ways in which hydrogen affects the deformation and fracture processes in amorphous alloys might be different from those in crystalline alloys. The changes in mechanical properties and fracture mode due to HE in amorphous alloys have been well documented, l~-sj but Very few quantitative studies on the embrittlement have been performed. In some of our pre- vious work, it was observed that an Fe-Ni base amor- phous alloy Fe40Ni38Mo4Bt8 could be embrittled in ambient hydrogen gas 1141 and cracking could be induced by static charging, tjsj In this article, we studied further J.-J. LIN, formerly Graduate Student, Department of Materials Science and Engineering, National Tsing Hua University, is Research Scientist, Materials Research Laboratories, Industrial Technology Research Institute, Hsinchu, Taiwan. T.-P. PERNG, Professor and Chairman, is with the Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan. Manuscript submitted May 28, 1993. the effects of hydrogen charging on the mechanical prop- erties of this alloy. The effects were quantitatively cor- related with hydrogen diffusivity and concentration within the specimen. II. EXPERIMENTAL The material studied was the commercial metallic glass Fea0Ni3sMo4B18 (Metglas 2826MB). The as- received ribbon was approximately 25.4-mm wide and 25-~m thick. Reduced-section specimens, with the width and length of the gage section being 6.3 and 25 mm, respectively, were prepared. Prior to the test, they were polished with 600-grit emery paper to give a fresh surface, and the thickness was reduced to about 20/zm. The area outside the gage section was insulated with lacquer. Hydrogen charging at current densities of 0.2, 0.5, and 2.0 mA/cm 2 was carried out prior to and during deformation at room temperature (20 ~ using a plat- inum anode in a solution of 0.1 N H2SO 4 with an ad- dition of 5 mg/L NaAsO2 to promote the ingress of hydrogen. A tensile test was performed at room tem- perature at strain rates in the range 7.58 • 10 6 to 2.67 10 -4 s -1. Aluminum shims were glued onto both ends of the specimen for gripping. Three to five speci- mens in each condition were tested and the arithmetic mean was calculated. The fracture surfaces were ex- amined by scanning electron microscopy (SEM). III. RESULTS A. Tests in Air For the specimens tested in air, plastic deformation occurred simultaneously with fracture. The linear rela- tionship of stress vs strain extended to the fracture point. The ultimate tensile strength (UTS) was about 2100 MPa, and the fracture strain and Young's modulus were 1.6 pct and 130 GPa, respectively. These values METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JANUARY 1995 197

Embrittlement of amorphous Fe40Ni38Mo4B18 alloy by electrolytic hydrogen

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Embrittlement of Amorphous Fe40Ni38Mo4B18 Alloy by Electrolytic Hydrogen

J.-J. LIN and T.-P. PERNG

The effects of hydrogen on the tensile properties and fracture processes at room temperature were investigated. Specimens were tested at various strain rates in air or under different cathodic charging-current densities. The slopes of the stress-strain curves were essentially identical for all the specimens, except that the fracture points varied under different test conditions. Macro- scopically, hydrogen only affected the elastic deformation behavior, but microscopically, the embrittlement was caused by the heterogeneous nucleation of localized plastic deformation. The degree of hydrogen embrittlement increased as the charging current increased or as the strain rate decreased. With the same charging current and time, longer dynamic charging resulted in more severe embrittlement. Before fracture took place, the strength of the alloy could be com- pletely restored if hydrogen had been removed. Hydrogen diffusivity and solubility were used to draw the time-dependent hydrogen concentration profiles for the specimens under different charging conditions. The difference in the mechanical properties was correlated with the hy- drogen concentration within the specimen.

I. INTRODUCTION

IT has long been recognized that hydrogen may embrittle many crystalline metals, including Fe- and Ni-base alloys. The embrittlement is manifested by the nonductile fracture mode, reduced ductility, and reduced tensile strength in a tensile test. Recently, some Fe- and Ni-base amorphous alloys also have been found to ex- hibit similar hydrogen embrittlement (HE) in a tensile or bending test by cathodic charging in acidic solutions, f~-81

Amorphous alloys are distinguished from the crystal- line alloys, particularly in their deformation behav- ior. M9-~31 Macroscopically, they behave in a brittle manner. In a tensile test, the linear relationship of stress strain extends to the fracture point. Plastic deformation occurs only shortly before or simultaneously with frac- ture. Plastic strain is localized in sharp shear bands, c~~ Their fracture is preceded by large local plastic shear, which produces a featureless smooth zone, followed by a catastrophic shear, which produces a "vein" or "ridge" pattern. 1~~ The lack of ductility has been ascribed to the inhomogeneous deformation. Therefore, the ways in which hydrogen affects the deformation and fracture processes in amorphous alloys might be different from those in crystalline alloys.

The changes in mechanical properties and fracture mode due to HE in amorphous alloys have been well documented, l~-sj but Very few quantitative studies on the embrittlement have been performed. In some of our pre- vious work, it was observed that an Fe-Ni base amor- phous alloy Fe40Ni38Mo4Bt8 could be embrittled in ambient hydrogen gas 1141 and cracking could be induced by static charging, tjsj In this article, we studied further

J.-J. LIN, formerly Graduate Student, Department of Materials Science and Engineering, National Tsing Hua University, is Research Scientist, Materials Research Laboratories, Industrial Technology Research Institute, Hsinchu, Taiwan. T.-P. PERNG, Professor and Chairman, is with the Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan.

Manuscript submitted May 28, 1993.

the effects of hydrogen charging on the mechanical prop- erties of this alloy. The effects were quantitatively cor- related with hydrogen diffusivity and concentration within the specimen.

II. EXPERIMENTAL

The material studied was the commercial metallic glass Fea0Ni3sMo4B18 (Metglas 2826MB). The as- received ribbon was approximately 25.4-mm wide and 25-~m thick. Reduced-section specimens, with the width and length of the gage section being 6.3 and 25 mm, respectively, were prepared. Prior to the test, they were polished with 600-grit emery paper to give a fresh surface, and the thickness was reduced to about 20/zm. The area outside the gage section was insulated with lacquer.

Hydrogen charging at current densities of 0.2, 0.5, and 2.0 mA/cm 2 was carried out prior to and during deformation at room temperature (20 ~ using a plat- inum anode in a solution of 0.1 N H2SO 4 with an ad- dition of 5 m g / L NaAsO2 to promote the ingress of hydrogen. A tensile test was performed at room tem- perature at strain rates in the range 7.58 • 10 6 to 2.67 • 10 -4 s -1. Aluminum shims were glued onto both ends of the specimen for gripping. Three to five speci- mens in each condition were tested and the arithmetic mean was calculated. The fracture surfaces were ex- amined by scanning electron microscopy (SEM).

III. RESULTS

A. Tests in Air

For the specimens tested in air, plastic deformation occurred simultaneously with fracture. The linear rela- tionship of stress vs strain extended to the fracture point. The ultimate tensile strength (UTS) was about 2100 MPa, and the fracture strain and Young's modulus were 1.6 pct and 130 GPa, respectively. These values

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JANUARY 1995 197

are independent of the strain rate in the range studied (7.58 • 1 0 -6 t o 2.67 x 10 4 S-I) and comparable to those reported for iron- and iron-nickel-base amorphous alloys.t3,8.1 l]

B. Tests for Charged Specimens

Four categories of experiments were designed to ex- amine the features of hydrogen embrittlement in Fe4oNi38MoaBis. The results are given in Sections 1 through 4.

1. Effect of strain rate Straining was started after the specimens had been

precharged at 0.2 m A / c m 2 for 5 minutes. Five strain rates (from 7.58 x 10 -6 t o 2.67 x 1 0 -4 S -l) were se- lected. Embrittlement was observed in all cases, except for those tested at the fastest rate. The degree of em- brittlement increased as the strain rate decreased or as the charging time increased. When tested at the slowest strain rate, 7.58 x 10 6 s-l , the UTS was reduced to 1190 MPa, i.e., 43 pct less than that in air. The UTS measurements of these specimens tested under various strain rates are presented based on the total charging time, as illustrated by the curve in Figure 1.

2. Effect of charging current Straining was started after the specimens had been

precharged at various current densities (0.2, 0.5, and 2.0 m A / c m 2) for 5 minutes. The strain rate was fixed at 7.58 x 10 -6 s - j . The degree of embrittlement in- creased as the current increased. Reductions of 58 and 75 pct in UTS were observed for 0.5 and 2.0 m A / c m : , respectively. Because they fractured at lower stresses, the total lengths of charging time were shorter. The UTS measurements for these specimens are also included in Figure 1.

O_ 3E

f--

200C'

1000

F%0 Ni38Moz. BlI~ Current Oensity(mA/cm 2) �9 : 0 . 2 o:05 ~:2

.

I I I t 1 10 20 30

Charging Time (min)

Fig. I - - U T S of Fe4oNi38Mo4B18 tested in 0.1 N H2SO4 + 5 m g / L NaAsO: under various charging currents and strain rates; T = 293 K.

3. Different combination of precharging and dynamic charging Specimens that had been precharged at 0.2 m A / c m 2

for various periods of time were tested in the following conditions:

(a) As described previously for specimens precharged for 5 minutes and then tested at 7.58 x 10-6s -1 under dynamic charging, the UTS was reduced by 43 pct and the total charging time was 25 minutes. (b) Straining was applied after 24 minutes of pre- charging. The strain rate was set at 2.67 x 1 0 -4 s -1

under dynamic charging. The specimens fractured within 1 minute. Therefore, the total charging time was also about 25 minutes, but the UTS was reduced only by 30 pct. (c) Specimens were tensile tested in air at a rate of 1.33 x 10 -5 s -1 after having been precharged for 30 minutes. It took about 15 minutes for them to frac- ture. The UTS was reduced by 25 pct.

4. Reversibility of hydrogen effect Two test sequences were performed to examine if the

hydrogen effects were reversible:

(a) Specimens were precharged at 0.2 m A / c m z for 30 minutes or at 2.0 m A / c m ~ f o r 20 minutes, baked at 120 ~ for 1 hour, and then tested in air at a strain rate o f 7 . 5 8 X 10 -6 S I. The UTS measurements were the same as those for uncharged specimens: 2100 MPa. (b) Specimens were precharged at 0.2 m A / c m 2 for 5 minutes, tensile tested at 7.58 x 1 0 -6 s - l under dy- namic charging to a stress of 1150 MPa (this is 97 pct of 1190 MPa), and held in air at room temperature for 24 hours. The test with the same rate was then resumed in air. The UTS was recovered to that of uncharged specimens, i.e., 2100 MPa.

C. Fracture Features

The slopes of the stress-strain curves were essentially unchanged for all specimens, except that the fracture points varied under different test conditions. However, the fracture path in the gage section was 45 deg to the tensile axis when tested in air and changed to 90 deg for hydrogen-charged specimens. For uncharged specimens fractured in air, the fracture surface consisted of a smooth region produced by local plastic shearing and a veinlike region produced by plastic instability, as seen in Figure 2(a). For hydrogen-charged specimens, most regions of the fracture surface were similar to those for uncharged ones, except they exhibited some cellular pat- tern. In general, only two or three cellular regions, each ranging from 200 to 700 /xm in length and extending through the whole thickness of the specimen (20 /zm), were observed for each specimen. The width of the gage section was 6.3 mm, and the fraction of cellular regions was only about 15 to 30 pct. These areas were randomly placed on the surface, and the fraction seemed to be in- sensitive to the charging time. Typical cellular patterns from different charging currents are shown in Figures 2(b) and (c). For the specimens whose strengths were completely recovered because of baking, the fractographic feature was also recovered to that of the uncharged specimens (Fig. 2(d)).

198--VOLUME 26A, JANUARY 1995 METALLURGICAL AND MATERIALS TRANSACTIONS A

(a) (b)

(c) (d)

Fig. 2 - - S E M fractographs for Fe4oNi38Mo4B~8: (a) air; (b) 0 .2 m A / cm 2, ~ = 7.58 x 10 6 s ~; (c) 2 . 0 m A / c m 2, f = 7.58 x 10 - 6 s ~; and (d) degassed.

IV. DISCUSSION

As shown in Figure 1, the degree of hydrogen embrittlement of amorphous Fe40Ni38Mo4B18 alloy by ca- thodic charging increased with the charging time after an "incubation time" until saturating at a limiting level. Severe embrittlement occurred after hydrogen had pen- etrated deeply into the specimen. If the charging time was too short, hydrogen distributed only within the sur- face layer, and the reduction in strength was not signif- icant. When the charging-current density was raised, the embrittlement also increased. This could be ascribed to the higher concentration of hydrogen associated with the higher charging current. If the hydrogen-charged spec- imens had been outgassed completely, not only the me- chanical strength but also the fractographic feature were recovered to those of the uncharged specimens. Even for the specimens tensile tested under dynamic charging to 97 pct of the reduced UTS, the effect of hydrogen was still reversible. In addition, we found that this alloy frac- tured without macroscopic plastic deformation when tested either in air or under dynamic charging. Macro- scopically, hydrogen seemed to affect the alloy only in the elastic region, and fracture occurred instantly when a critical condition had been reached. No shear or micro- crack was induced before fracture took place. A similar reversible effect of hydrogen on the deformation process has been reported previously, t2,61

Several mechanisms have been proposed to explain hydrogen embrittlement in crystalline materials, e . g . , hydrogen may (a) reduce the lattice cohesive f o r c e , [j6'17]

(b) reduce the surface energy, l~Sj (c) interact with dis- locations, 119j or (d) build up an internal pressure. ~2~ Some of these mechanisms have also been applied to explain the effect in amorphous alloys. For instance,

Ashok et al. tSI examined the fractographic features of alloys embrittled by liquid metal and hydrogen and pro- posed that the effect was through the adsorption-induced plasticity ( i . e . , enhanced shear at crack tips). On the other hand, Schroeder and K6ster t61 explained the em- brittlement behavior by assuming that hydrogen reduced the cohesion between interatomic bonds and/or the quenched-in excess free volume of the metallic glass. In a previous article, we found that F e 4 0 N i 3 8 M o 4 B 1 8 could be cracked simply by prolonged static charging with hy- drogen. 115~ The cracking was ascribed to an internal hy- drogen pressure built up around heterogeneous sites. The critical pressures for crack initiation under different charging conditions could be calculated based on the sur- face hydrogen concentration and hydrogen diffusivity data. The effect of internal hydrogen pressure was less important in this study, because the durations of the tests were too short to allow the build up. Determining whether the embrittlement observed here was caused by enhanced shear at crack tips or by decohesion between the atomic bonds is not the aim of this study. Either of these two theories, however, necessitates a higher hy- drogen concentration for a higher degree of embrittlement.

When the results of experiments C(1) and C(2) were compared, it was seen that with the same charging cur- rent and time, C(1) resulted in more severe embrittle- ment than C(2). This larger effect could be ascribed to the larger hydrogen flux diffusing into the specimen under the C(I) test condition. In the C(1) test, the spec- imen was subjected to a much longer time of dynamic straining, i . e . , 20 minutes compared with 1 minute for the C(2) test. During this period of time, more hydrogen was accumulated inside the specimen. Because the flux is determined by the product of diffusivity and surface concentration of hydrogen, it indicates that either dif- fusivity or surface hydrogen concentration increases under dynamic straining. Zakroczymski found that for iron, elastic deformation resulted in a slight increase in the permeation rate but with no change in diffusivity, t211 which was ascribed to an increased hydrogen solubility in the elastically expanded lattice. Frankel and Latanision proposed that increased input hydrogen con- centrations were associated with slow strain rates for nickel. 1221 Because hydrogen transport behavior in this alloy is similar to that in fcc alloys, ~23,~-41 it is suspected that there is also some change in hydrogen flux resulting from elastic straining in this alloy. More detailed ex- periments should be performed to clarify this.

To model how the hydrogen concentration affects the severity of embrittlement of this sample, a simple anal- ysis is made. The hydrogen concentration profiles for the specimens under various charging conditions are dis- played in Figure 3. Hydrogen diffusivity is estimated from a 2previous permeation experiment (D ~ 7 • 10 -15 m /s), I231 and the surface concentration before straining Co is extrapolated from the data in a previous work ILSj to be 3.8 x 10 -2 atom H / M (5.1 x 10 3 mol H/m3). In Figure 3 (a), a test from category 1 is illus- trated. The charging time was short (about 6 minutes), hydrogen penetration depth was only 3 to 4 /zm, and the concentration of hydrogen was too low to affect the strength of this alloy. If the charging was prolonged

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 26A, JANUARY 1995-- 199

g

w e - o

g

P 5min

+ D l m i n

zxUTS:O

[--Thickness (20gm)~

(a)

C.

P 5min

+ D 20min

', , ' .UTS= 4.3% i \ /

,,7 "% ,2'

(b)

P 2 4 m i n

+ D 1 m i n

C.

P 3 0 m i n

+ O.G. 15min

�9 " U ' fS = 25*/ , .

(c) (d)

C.

C.

Fig. 3 - -Schemat i c hydrogen concentration profiles under various charging conditions and the associated reductions in UTS. Charging current: 0.2 m A / c m 2. P: precharge; D: dynamic charge; O.G.: outgas; and Co: surface concentration before straining (= 3.8 x 10 2 atom H/M). The dashed curves in (a) through (c) are the modified profiles when the specimens were subjected to straining.

(e.g., experiments C(1) and C(2)), hydrogen penetrated more deeply, as shown in Figure 3(b) and (c), respec- tively. The solid curves represent the concentration pro- files for the specimens in the absence of straining. Based on Frankel and Latanision's proposal, 122] however, the surface concentration was slightly raised because of the straining. The concentration profiles should be modified. For experiment C(1), the straining duration was much longer and the concentration increment extended to the middle of the specimen, as illustrated by the dashed curve in Figure 3(b), whereas for experiment C(2), the concentration increment occurred only near the surface (Figure 3(c)). The magnitude of increase in surface con- centration is difficult to determine, and therefore the dashed curves are only schematic in nature. The differ- ence in the concentration between Figures 3(b) and (c) accounts for the difference in hydrogen effect on UTS (i.e., 43 vs 30 pct). Similar modification on the con- centration profile is also made for Figure 3(a).

In experiment C(3), the specimen had been pre- charged for 30 minutes before it was tested in air. Some of the dissolved hydrogen diffused out during the test. The total amount of hydrogen within the specimen de- creased after 15 minutes of outgassing, but the distri- bution of hydrogen became more uniform, except that the concentration of hydrogen near the surface dropped to zero, as shown in Figure 3(d). The UTS was reduced by 25 pct. Compared with Figure 3(a), it seems that only the hydrogen atoms that have penetrated deep enough would contribute to the embrittlement. According to Figure 3(d), with Co being 3.8 • 10 2 atom H/M, the average hydrogen concentration within the specimen is calculated to be about 1.3 x 10 2 atom H/M. This con- centration is several orders of magnitude larger than

those required for austenitic stainless steel to yield a sim- ilar degree of embrittlement) 25,26J

The ability for the near-surface layer to accommodate hydrogen was limited in nature. When a higher charging current was applied, (e.g., 0.5 or 2.0 mA/cm2), a higher hydrogen concentration was built up. Greater reductions in UTS were observed, as shown in Figure 1. On the other hand, from the concentration profiles shown in Figure 3(b) and (d), it is seen that a higher average con- centration in the middle does not necessarily lead to a higher reduction in UTS. The higher concentration of hydrogen near the surface still plays an important role in the failure process of the alloy. A quantitative expla- nation for the embrittlement effect has to take more fac- tors into account, e.g. , interaction between hydrogen and the metal atoms, the different deformation modes in the near-surface layer and in the middle, and the fracture initiation process.

The fractographic features of the charged specimens exhibit a cellular pattern, which can be considered as a very fine-scaled dimple structure. It is an indication of heterogeneous nucleation of localized plastic deforma- tion. [2] The smaller cell size observed for 2.0 mA/cm 2, as compared in Figures 2(b) and (c), indicates that more nucleation sites for fracture are associated with the higher charging current. Because fractographic features are reversible as the strength of specimen is recovered by outgassing, this localized plastic deformation proba- bly occurs only shortly before or simultaneously with fracture in a tensile test.

Finally, it is worth emphasizing that the region exhib- iting the cellular pattern constituted only a small fraction of the entire fracture surface. Most regions were still similar to that for an uncharged one. This implies that hydrogen embrittlement takes place preferentially in some specific regions or, in other words, in some re- gions of higher heterogeneity. It is conceivable that al- though amorphous alloys do not have defects, such as a grain boundary or dislocation, heterogeneous regions are induced intrinsically by rapid quenching.

V. CONCLUSIONS

1. Amorphous Fe40Ni3sMo4Bi8 alloy could be embrittled by cathodic charging with hydrogen. The degree of embrittlement increased as the charging current in- creased or as the strain rate decreased. With an av- erage uniform concentration of about 1.3 x 10 -2 atom H / M within the alloy, the UTS was reduced by 25 pct.

2. With the same charging current and time, longer dy- namic charging resulted in more severe embrittlement.

3. Macroscopically, hydrogen only affected the elastic deformation behavior, but microscopically, the em- brittlement was caused by heterogeneous nucleation of localized plastic deformation. The effect was re- versible. The UTS of the alloy was completely re- stored by degassing of hydrogen before fracture took place.

200--VOLUME 26A. JANUARY 1995 METALLURGICAL AND MATERIALS TRANSACTIONS A

ACKNOWLEDGMENTS

This research was supported by the National Science Council of the Republic of China under Contract No. NSC 80-0405-E-007-11. The specimen material was supplied by the Materials Research Laboratories of ITRI.

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