JMEPEG (1997) 6:534-544 �9 International

The Susceptibility of 90Cu-10Ni Alloy to Stress Corrosion Cracking in

Seawater Polluted by Sulfide Ions A. El Domiaty and J.N. Alhajji

Electrochemical polarization measurements and slow strain rate tests (SSRT) of a 90Cu-10Ni alloy in highly sulfide polluted seawater were conducted to investigate stress-corrosion cracking (SCC) behav- ior. The severity of the SCC depends on the sulfide concentration in the seawater. The severity increases as the concentration increases. Because the major time in SCC is spent in the initiation process of the propagating crack, the fracture toughness has only a minor effect in the component life failed by SCC. The SCC behavior of CDA706 is strictly linked to sulfide concentration in the range of 100 to 1000 ppm. The general corrosion of Cu-Ni alloys in low (<100 ppm) and high (>100 ppm) sulfide polluted seawater increases due to the selective copper dissolution. Cyclic polarization measurements confirmed that the corrosion rate decreases slightly as the sulfide concentration increases. Pitting tendency was high in the low concentration range of sulfide and low in the high concentration range. The presence of stresses in SCC removes the protective layer as it increases during testing of the specimen or during the actual ser- vice of a component. The authors propose that film rupture occurred, and two proposed SCC mecha- nisms were operational, namely sulfide stress cracking associated with the anodic dissolution in the low sulfide concentrat ion range and hydrogen embritt lement, which was dominant in the high sulfide concentration range. It was found that a synergism exists between sulfide and stress that enhances the effect of the latter.

Keywords CDA706, copper-nickel alloy, corrosion, crack initiation, crack propagation, fracture toughness, intergranular corrosion cracks, open circuit potential, polarization, seawater, slow strain rate tests, stress-corrosion cracking, sulfide

1. Introduction

Copper alloys are used extensively in water distribution systems, water treatment units, and condensers and heat ex- changers, where fresh or salt water is used for cooling. These units are often the critical components for the processing plant in electricity generating, oil, and chemical industries. The use of copper alloys in any of these engineering applications is often limited by its susceptibility to stress-corrosion cracking (SCC). The selection of a specific alloy for a certain applica- tion is usually based on laboratory test results that were ob- tained in the simulated environment.

Three main methods for SCC testing include constant load, constant strain, and constant strain rate. Constant load and con- stant strain rate techniques are more discriminating but gener- ally require large loading frames. Conversely, constant strain tests require only small loading devices, and a large number of specimens can be tested simultaneously under different stress levels. The criterion for SCC resistance in this type of test is the time to initiate and propagate a micro crack, which can be seen at low power magnification, not the time to failure, because the specimens do not fail. Because each testing technique has ad- vantages and disadvantages, there is a need for a simple and re- liable SCC test method. A simple automated stress corrosion

A. E! Domiaty and J.N. Alhajji, Kuwait University, Safat 13060, Kuwait.

testing method called the automated stress corrosion ring (ASCOR) test was developed by Schra and Groep (Ref 1) to test aluminum alloys according to ASTM G 44. A specific load decrease (2%) is used as the criterion for SCC initiation. The main advantages of the ASCOR test method are that a large number of specimens can be tested simultaneously, and the SCC initiation lines can be determined according to a clearly defined criterion without time consuming and subjective in- spections and without disturbing the test procedure. The slow strain rate test (SSRT) uses (Ref 2) tensile test specimens mounted in stiff-frame machines and strained at the rate of 10 -7 to 10 -5 s -1 in the presence of a specific environment. Strain rates in this range promote SCC, but the absence of cracking is no ensurance of immunity to SCC. Various methods are used to assess the results when SCC is observed (Ref 3). These include the area under the stress-strain curves, time to failure, crack ve- locity, and the ratio of fracture stress in a medium to fracture stress in air.

1.1 Slow Strain Rate Testing

Because of the distinct advantage of rapidity of the slow strain testing (SST) technique in the assessment of SCC sus- ceptibility, the present investigation employed SST for the SCC studies on copper-nickel alloys in a sulfide polluted envi- ronment. I f the strain rate is too high, ductile fracture occurs be- fore the necessary corrosion reactions take place. However, at too low a strain rate, corrosion can be prevented by repassiva- tion or film repair so that the necessary reactions of bare metal cannot be sustained. The optimal strain rate for testing the SCC susceptibility depends primarily on the metal/environment system or the crack propagation rate (CPR). Usually for a sys- tem in which the CPR is small, a slow strain rate is required, and in cases where the CPR is high, a relatively fast rate is suit-

534---Volume 6(4) August 1997 Journal of Materials Engineering and Performance

able. For admiralty brass in sulfite or nitrate solution, the CPR is approximately 10 -7 m/s for which a strain rate of 10 -5 s -1 is adequate. In the late 1960s, Parkins (Ref 3) described a test arrangement where an immersed tension specimen in a cor- rosion solution polluted at a slow constant extension rate. The data obtained from SSTs provide corrosion engineers with a greater degree of confidence in choosing compatible material/environment combinations. Not only can this method satisfy pass/fail criteria, but it can also provide data to judge the relative degree of SCC susceptibility among test variables, conditions that cause SCC to result in a lower maximum load and a loss of ductility. Therefore, by comparing maximum load, reduction in area or elongation and noting the degree of secondary cracking, a particular variable can be judged to re- tard or accelerate cracking.

In addition to the testing methods for SCC, attention should be focused on the corrosion/film formation tendency of un- stressed specimens in a particular environment. The corrosion behavior of copper-nickel alloys in unpolluted and sulfide pol- luted seawater was investigated by Eiselstein et al. (Ref 4). They found that the copresence of sulfide and oxygen in seawa- ter results in corrosion rates much higher than those measured when the alloy is exposed alternately to sulfide and oxygen. In aerated sulfide polluted seawater, accelerated corrosion ap- pears to be the result of the sulfide preventing the formation of a protective oxide corrosion product layer. The behavior of 70Cu-30Ni alloy in quiescent sulfide polluted and unpolluted seawater at 25 ~ was examined by Beccaria et al. (Ref 5). They demonstrated that the presence of sulfides enhances the average corrosion rate in a range of concentration (0 to 10 ppm) while promoting preferential copper dissolution. In deaerated seawater, corrosion behavior was found to be strictly linked to sulfide concentration. In low polluted seawater, the mean cor- rosion rate of the alloy increases, and slightly selective dissolu- tion occurs. In highly polluted sea water, the corrosion rate decreases while slightly selective nickel dissolution occurs. The corrosion of copper-nickel alloys in sulfide-polluted seawater was also studied by Alhajji and Reda (Ref 6-11). They found that sulfide concentration plays a major role in the corro- sion process. They concluded that sulfide acts as a catalyst, and the copper sulfide is the first compound to precipitate in sulfide polluted seawater.

Most of the early information regarding the SCC of copper alloys was based on service experience. Such information is useful but qualitative, and the environmental constituents or conditions that led to the cracking are unknown. In the past sev- eral decades, the study of SCC was greatly accelerated, and the causes and the mechanisms for the cracking behavior were ad- dressed. The SCC of copper-nickel alloys was studied by Thompson (Ref 12), and he found that the alloys are practically immune to SCC in seawater. Popplewell and Gearing (Ref 13) showed that only a few of these 25 copper-base alloys are sus- ceptible to SCC in industrial and marine environments. Using the SST technique, Habib and Husain (Ref 14) concluded that the most severe SCC was observed on the 70Cu-30Ni specimen in seawater with 3120 ppm sulfide at 25 ~ It was also con- cluded that the 90Cu-10Ni and 70Cu-30Ni alloys are suscepti- ble to SCC in sulfide seawater, which was based on the ratio of ultimate tensile strength, elongation percent, and time to fail-

ure of samples tested in solutions compared to those tested in air at moderate temperature.

In this study, an experimental program was conducted using the SST technique to determine the susceptibility of Cu-Ni al- loys in sulfide polluted seawater. This experimental program included a scanning electron microscope (SEM) study of the fracture surface as an important assessment tool that was missed by some investigators (Ref 14) to determine the type of fracture and gain an insight into fracture mechanism. Also, the potential drop technique is used during the test to determine the onset of crack initiation. In addition, various electrochemical tests were conducted under unstrained conditions to elucidate the mechanism responsible for the appreciable increase in cor- rosion tendency and the embrittling effect on 90Cu- 10Ni in the presence of sulfide.

2. Experimental Procedures

2.1 Materials

The material used for this study was 90Cu-10Ni (CDA706) in the form of tube, as received from the supplier. The inner and outer diameters of the tube were 15.5 and 19 mm, respectively. The mechanical properties of the materials were determined by the standard tensile test performed according to ASTM E 8-83. Flat specimens of 15 mm diameter and 2.5 mm thickness were subjected to various electrochemical tests under unstrained conditions to evaluate the corrosion and pitting tendency of the specimen under the various environmental conditions. Figure 1 shows the corrosion test specimen dimensions. Figure 2 shows the tensile test specimen mounted in the corrosion cell.

I I

2.2 SCC Test Cell and Procedures

Figure 2 represents a schematic drawing of the entire strain- ing electrode/corrosion cell assembly. The corrosive cell is a plexy glass tube with a rubber cork in the lower end and a plexy glass plate on top. The specimen in the form of a pipe was mounted (Fig. 2) and subjected to uniaxial loading by the In- stron testing machine. The cell was designed so that the gage length of the specimen is surrounded by the corrosive environ- ment. The counter electrode was in the form of cylindrical platinum mesh surrounding the entire gage length of the speci- men and reference electrode. The cell was connected to the po- tentiostat to monitor the potential change during the SSRT. Testing was limited to aqueous solutions at room temperature

l j LAPPED SURFACE 8 - 10 RM$ One Side Only

2.5 mm

Fig. 1 A schematic diagram of the test specimen used in the corrosion tendency evaluation

Journal of Materials Engineering and Performance Volume 6(4) August 1997--535

in different concentrations of sulfide in seawater at 0, 100, 300, 500, and 1000 ppm.

2.3 Electrochemical Tests

Samples of CDA706 were prepared by a wet grinding proc- ess with 240-grit SiC paper and wet polished with 600-grit SiC paper until previous coarse scratches were removed. Then the specimens were rinsed and dried. These procedures were com- pleted one hour prior to the experiment (Ref 8). Corrosion measurement experiments were conducted using synthetic seawater. This was prepared with distilled water and standard seawater salt. Standard seawater salt (Marinemix and Bio-Ele- ments from Wiegandt GMBH & Co., F. R. Germany) was used to reduce the variability effects resulting from conducting measurements using natural seawater. The sulfide was intro- duced in the seawater using research grade sodium sulfide (Na2S). The level of sulfide in the seawater was checked by the iodometric method of analysis (Ref 9). Analytical grade chemicals were used in all experiments.

Electrochemical corrosion measurements were taken at 20 ~ for all previously mentioned conditions using a computer controlled potentiostat/galvanostat (EG&G; Z.I. Petite Mon- tagne Sud, France). A modified electrochemical corrosion test cell was used, in which a large cylindrical platinum mesh counter electrode was incorporated to obtain a uniform electri- cal field among the working, reference, and counter electrodes. Preliminary corrosion measurements were always performed for 4 h using small amplitude cyclic voltammetry (SACV) over the range of +_5 mV from Et= 0 as previously explained (Ref 9). In addition, cyclic potentiodynamic scans were run after 4 h to gain an insight into the corrosion mechanism in the presence of a sulfide pollutant. The electrochemical tests were run with a saturated Calomel electrode (SCE). The stability of the SCE was checked following each experiment against a fresh SCE to ensure the reliability of the experiments. The scan rate of the experiments was 0.166 mV/s. The experiments were con- ducted repeatedly with fresh samples under the same environ- mental conditions to obtain reproducible results. Generally, reproduced results were excellent.

3. Results and Discussion

An electrochemical study for the unstressed 90Cu-10Ni al- loy specimens in unpolluted and sulfide polluted seawater was completed prior to the assessment of SCC susceptibility of the alloy in these environments. It is known (Ref 6-11, 15-21) that sulfide results in a serious deterioration of Cu-Ni alloys be- cause of the film formation/breakdown tendency of cop- per/nickel corrosion in sulfide environments. The following mechanisms are proposed to counteract the effects of sulfide on this alloy.

3.1 Proposed Corrosion Mechanism of Cu/Ni Alloys in Seawater

Generally, in sulfide polluted seawater, the following se- quence of reactions are possible :

0 2 + 4H20 + 4e- --> OH-

Cu + HS- --> Cu (HS-)ad s

Ni + HS- --> Ni (HS-)ads

Cu (HS-)ad s--~ Cu (HS) + e-

Ni (HS-)ads---> Ni (HS) + + 2e-

Cu (HS) ~ Cu + + HS -

Ni (HS) + --~ Ni 2+ + HS -

2Cu + + HS- + OH-----> Cu2S$ + H20

Ni 2+ + HS- + OH- ----> NiS$ + H20

(cathodic reaction) (1)

(2a)

(2b)

(3a)

(3b)

(4a)

(4b)

(5a)

(5b)

Once the cuprous-sulfide and nickel-sulfide are formed, it is concluded that the intensification of corrosion is not directly linked with the dissolved sulfide in the electrolyte but rather with a secondary effect caused by the action of Cu2S or NiS, where the sulfides act as effective cathodes (Ref 22). The ad- sorption of HS- ions (reactions 2a,b) creates negatiye potential that increases the speed of the hydrogen discharge reaction through another simultaneous cathodic reaction as follows (Ref 22):

HS- + e- --+ Had s + S = (6)

Had s + Had s ---> H 21" (7)

where S = at the pH of the solution (9.7 to 9.75) is not stable, and thus the following reaction takes place:

S = + H20 ~ HS- + OH- (8)

and thus, hydroxide ions are produced resulting in a localized increase in the pH of the solution and in the formation of hy- droxides that occur through the followingreactions:

Cu + + OH- ~ Cu (OH),I, (9a)

Ni 2+ + 2OH- ---> Ni (OH)2,1, (9b)

The formation of sulfides occurs through the following reac- tions:

2Cu + + S = --> Cu2S,l, (lOa)

Ni 2+ + S = ----> NiS,[, (10b)

536---Volume 6(4) August 1997 Journal of Materials Engineering and Performance

Thus, corrosion products are brittle, nonadherent, and consist of a nonprotective mixture of hydroxides and sulfides of cop- per and nickel. The hydrogen produced by reactions 6 and 7 is adsorbed on the surface, and the specimen becomes susceptible to hydrogen cracking and enhances the brittling effect of sul- fide. This is particularly important in this study because the tests were performed at SST (Ref 23), and the film produced was nonprotective.

Figure 3 shows the corrosion potentials of CDA706 in sul- fide-polluted seawater. It shows that increasing sulfide levels result in a shift in corrosion potential in the active direction.

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Fig. 3 Open circuit potential diagram for unstrained 90Cu- 10Ni alloy in sulfide polluted seawater

This shift in potential is directly related to the reaction occur- ring on the surface as outlined earlier. Clearly, increasing sul- fide concentration results in the most active shift in potential, and hydrogen evolution then becomes a possible cathodic par- tial process. In addition to reactions 6 and 7, another reaction becomes significant (Ref 24):

2H20 + 2e- --~ H 2 1" + 2OH- (11)

This reaction further supports the earlier supposition that sulfide also plays an indirect role in inducing hydrogen crack- ing. It is proposed that film rupture occurred, and two major SCC mechanisms were operational, namely sulfide stress cracking associated with the anodic dissolution and hydrogen embrittlement. Figure 4 shows that increasing the sulfide con-

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log [Sulfide Concentration, (ppm)]

The corrosion tendency of the 90Cu-10Ni alloy (CDA706) in aerated sulfide polluted seawater. (a) Linear plot. (b) Semi-log plot.

Journal of Materials Engineering and Performance Volume 6(4) August 1997--537

centration results in a sharp increase in the corrosion tendency of the Cu-Ni alloy under investigation. It should be emphasized that these tests were performed over a relatively short duration, and a significant corrosion product layer did not form that could lead to a leveling off or decrease in the corrosion rate over time due to diffusion limitations of the sulfide ions.

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Fig. 5 Cyclic potentiodynamic diagrams for the unstrained 90Cu- I 0Ni alloy in seawater

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The present research used very high sulfide concentrations ranging from 100 to 1000 ppm in seawater as the primary cor- roding medium for this SCC study. The polarization diagram (Fig. 5, 6) shows that the increase in potential in the anodic di- rection yielded an increase in current in the region close to the equilibrium potential. After the initial active dissolution, re- duction in anodic current values was noted for further increase in potential and remained little affected by the potential o f - 300 mV versus SCE for all sulfide polluted systems. It was re- ported (Ref 9) that the cuprous-oxide films formed during the initial exposure to unpolluted seawater were relatively smooth and nonporous, while the sulfur containing films, presumed to be stoichiometric and substoichiometric forms of cuprous-sul- fide, were relatively thick and porous. It was further stated that the presence o f dissolved sulfide or sulfide oxidation prod- ucts does not lead directly to accelerated corrosion, but rather that the porous cuprous-sulfide corrosion product formed in the polluted water interferes with the normal growth of the protective oxide films on subsequent exposure to unpolluted seawater.

The present research studies were performed under stagnant conditions; the temporary passive tendency of the corrosion product layer formed during the anodic polarization was noted for the concentrations of sulfide and the decrease in the in- crease in the sulfide concentration. This behavior was observed up to the potential value of approximately -300 mV versus

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Cyclic potentiodynamic diagrams for the unstrained 90Cu-10Ni alloy in sulfide polluted seawater

538---Volume 6(4) August 1997 Journal of Materials Engineering and Performance

SCE for all cases. However, increases in potential beyond this value of -300 mV, close to the equilibrium potential of the 90Cu-10Ni alloy in an unpolluted system resulted in the active

dissolution, and the behavior was the same for all sulfide pol- luted systems, irrespective of the concentration.

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Fig. 7 Stress-strain diagram for the 90Cu-10Ni alloy in air and sulfide polluted seawater during slow strain rate testing (SSRT)

3.1 Mechanical Results

Figure 7 shows the engineering stress-engineering strain curves for the 90Cu-10Ni alloy in the various environments. Table 1 gives the mechanical properties obtained from the standard tension test at room temperature. Typically, SCC se- verity is expressed as the ratio of a parameter from the SCC curve to the same parameters of the baseline curve. As this ratio decreases in unity, the SCC severity increases. Several parame- ters that are available for comparison include maximum load, total elongation, area under the curve, and failure load. The true

Table 1 Mechanical proper t ies of the 90Cu-10Ni alloy f rom the s t anda rd tension test a t room tempera tu re

Alloy property Cu/Ni:90/10

Tensile strength, MPa 350 Yield strength (0.2% offset), MPa 180 Elongation, % 36 Reduction in area, % 63

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Fig. 8 The embrittling effect of sulfide concentration during slow strain rate testing (SSRT) for the 90Cu-10Ni alloy in seawater. (a) Per- cent elongation. (b) Percent elongation (semi-log plot). (c) True fracture strain. (d) True fracture strain (semi-log plot)

Journal of Materials Engineering and Performance Volume 6(4) August 1997--539

fracture stress and true fracture strain can also be used to evalu- ate SCC severity.

Measurements of the reduction in area and elongation per- cent are indicators of specimen ductility loss in SSRT, which provides a convenient parameter for evaluating SCC severity. Figure 7 shows the stress versus strain curve resulting from SSRT for the CDA706 alloy exhibiting SCC. Ductility loss can also result from causes other than SCC (Ref 25), for example, hydrogen charging through Reactions 6, 7, and I 1 on the sur- face of the Cu-Ni alloy. In this study, where hydrogen coexists through Reactions 6, 7, and 11 with the causative agents of SCC (sulfide ions and oxidation products, e.g., HS-), interpre- tation of results must be made carefully. Table 2 shows the em-

90 , . , I , , . I , , . I , . , I , . . I,

~" 80- t . . ,

o

70-

, o

0 200 400 600 800 I000

Sulfide Concentration (ppm)

Fig. 9 The embrittling effect of sulfide concentration during slow strain rate testing (SSRT) for the 90Cu-10Ni alloy in seawater. (a) Time to failure (linear plot). (b) Time to failure

1

1200

-0.2

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~_., -0.4

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T i m e (106 S)

Fig. 10 Open circuit potential diagram for the 90Cu-10Ni al- loy in sulfide polluted seawater during slow strain rate testing (SSRT)

I I I I I

0.1 0.2 0.3 0.4 0.5

brittling effect of increasing the sulfide concentration on the re- duction in toughness of the Cu-Ni alloy measured by the area under the 6-~ curve. Table 2 reports stereo microscopic obser- vations and shows a general increase in the density and distri- bution of secondary stress corrosion cracks with an increase in sulfide concentration. Figures 8(a) and (b) show the embrit- tling effect of sulfide concentration during SSRT for 90Cu- 10Ni in seawater. Clearly, high levels of sulfide pollutant in seawater results in a significant ductility loss and reduction in fracture strain. The ductility given by the elongation percent was reduced from 30% when the concentration was 100 ppm to 15% when the concentration increased to 1000 ppm. The true fracture strain is a sensitive measure for SCC susceptibility. Figures 8(c) and (d) show the true fracture strain versus the sul- fide concentration.

Specimen elongation is directly related to the time of failure because a constant cross head extension is imposed during the slow strain rate test, and the load decreases and drops to zero at failure. Figure 9 shows a sharp decrease in the time to failure with increasing sulfide concentration. This further substanti- ates the embrittling effects of sulfide ion and its byproducts.

From the previous results, it can be concluded that in slow strain tests on 90Cu-10Ni in sulfide polluted seawater, the frac- ture strain and the time to failure provide a strict criterion for SCC susceptibility. The strain prior to crack initiation is also determined by electrochemical monitoring of the specimen during the test.

Figure 10 shows the electrochemical monitoring of the open circuit potential, El= o. It was found that a similar trend exists in this case when compared to the unstrained conditions shown earlier in Fig. 3. A sharp systematic potential shift in the active direction was observed in the strained conditions with increas- ing sulfide concentration. This shift in potential is directly re- lated to the reaction occurring on the surface of the alloy as outlined earlier. This active shift in potential results in hydro- gen evolution via Reactions 6, 7, and 11, which become a pos- sible cathodic partial process supporting the earlier predictions that the sulfide also plays an indirect role in inducing hydrogen cracking. It appears that this effect becomes more pronounced with increasing sulfide levels. The time dependent effects of sulfide on potential and mode of failure are broken down into two opposing effects: high levels of sulfide (_>300 ppm) and low levels of sulfide (_<100 ppm).

In the presence of high concentrations of sulfide (>300 ppm), the potential continues to increase from its initial active values, which indicates a mild increase in crack velocity. Early crack initiation and propagation causes a shift in potential in the noble direction indicative of a mild increase in propagating velocity of the crack. At high levels of sulfide, this effect is in- dependent of its concentration. The premature failure of the specimen in these environments can easily be attributed to the significant increase in hydrogen charging and the associated embrittling effect. This effect overwhelms the direct sulfide stress cracking contributed by sulfide and its byproducts. Thus, for high levels of sulfide, the nature of the fracture is dominated by hydrogen embrittlement, and the role of sulfide stress crack- ing is minimized.

For environments of low sulfide pollutant (_<100 ppm), an initial shift in the active direction is directly related to material

54if--Volume 6(4) August 1997 Journal of Materials Engineering and Performance

response to sulfide and, at later stages, crack initiation fol- lowed by crack propagation occurs. The contribution of hydro- gen in this case is minimal, and significant crack propagation occurs through the action of sulfide and its oxidation products. The initial continued potential drop appears to be the result of a film rupture process and the creation of an active surface at the slip steps. Figure 10 shows this type of behavior for the case of a 100 ppm sulfide polluted environment. A metastable re- gion of potential occurs at about 90 x 103 s, which can be ex- plained by pseudoequilibrium effects of the film free surface

Fig. 11 Longitudinal cross section of a tensile gage region showing secondary stress-corrosion crack for the 90Cu-10Ni alloy in 300 ppm sulfide (200x)

including the newly formed crack and the repassivation proc- ess due to sulfide and oxidation products. A large potential in- crease occurs, apparently due to crack formation in the final stage of the test. Figure 10 shows this type of potential rise. The electrode potential reaches a maximum, and the specimen breaks. It is expected that the fractured specimen will mainly show sulfide stress cracking, and the contribution of hydrogen is minimal. The total time to failure can also be determined from the data (Fig. 10) for the low sulfide levels, which are equal to 350 x 103 s (97 h). This is consistent with the failure time of the SST.

3.2 Morphology of Fracture

Metallography or fractography is always used to verify the presence of SCC after SSRT and can provide both quantitative and qualitative descriptions of SCC severity. Both optical and scanning electron microscopes were used to examine the failed specimens. The presence of secondary cracks on the surface of the tested material is a good indication of the severity of SCC. Table 2 shows the stereo microscopic observations with in-

Table 2 Toughness of the 90Cu-10Ni alloy measured by the area under the o-~ curve as a function of sulfide pollutant concentration

Sulfide environment, ppm Observations o-e area, M Jim 3

0 100

300

500

1000

Ductile material 106.19 Few secondary stress-corrosion 82.34

cracks in the necked region Increased secondary stress- 57.47

corrosion cracks in the necked region

Further increase in secondary 47.35 stress-corrosion cracks in the necked region

Secondary stress-corrosion 28.53 cracks over a large region of gage length

(a) (b)

Fig. 12 (a) A photomicrograph of the 90Cu-10Ni alloy in 500 ppm sulfide (1100x) showing the embrittled region of fracture. (b) A photo- micrograph of the 90Cu-10Ni alloy in 300 ppm sulfide (1500x) showing the ductile dimple type failure in the necked region

Journal of Materials Engineering and Performance Volume 6(4) August 1997--541

Fig. 13 A photomicrograph of the 90Cu-10Ni alloy in 300 ppm sulfide (700x) showing the intergranular crack propagation

Fig. 14 A photomicrograph of the 90Cu- 10Ni alloy in 500 ppm sulfide (7500x) showing the intergranular crack propagation region

(a) (b)

Fig. 15 transgranular cracking

creasing sulfide concentrations. Clearly, SCC of alloys is envi- ronment specific and is promoted by critical ranges of concen- trations of the culprit ions. Shallow penetrations under stress-corrosion circumstances indicate borderline conditions between regions of severe SCC and pitting corrosion.

Laboratory observations showed that increasing the con- centration of sulfide always increased the general corrosion and the deposits of corrosion byproducts. The specimen was encrusted with dark brown, loosely adhered scales. A range of features from severe general corrosion, pitting, and secondary stress corrosion cracks in the necked region were observed upon examining the longitudinal cross section of the tensile gage region (Fig 11). In general, secondary corrosion cracks on the surface were increased upon increasing the sulfide concen- tration, which increased the corrosivity of the environment (Table 2). Figures 12(a) and (b) show typical fractographic fea- tures of failed specimens.

A photomicrograph of the 90Cu-10Ni alloy in 500 ppm sulfide (4500x) showing the intergranular decohesive fracture with slight

The transition from brittle, intergranular propagation in the SCC region (Fig. 12a) to ductile, dimple propagation in the mechanical region at the bottom of the micrograph is shown in Fig. 12(b). The secondary cracks were completely inter- granular and oriented perpendicular to the direction of the ap- plied stress.

The features were typical of what would be expected in nor- mal SCC failure. Figure 13 shows extensive intergranular cracking where transgranular cracking was less prominent. The hydrogen ions resulting from Reactions 6, 7, and 11 are ad- sorbed and diffuse into the alloy, resulting in the production of localized cracking. The region of the alloy exposed at the crack tip as the crack propagates by virtue of hydrogen embrittlement and the applied stress are anodic to the oxidized sides of the crack and the adjacent surface of the material, which results in the continuation of the electrochemical attack and the further evolution and absorption of hydrogen. The triaxial state of

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(a) (b)

Fig. 16 A photomicrograph of the 90Cu- 10Ni alloy in 1000 ppm sulfide showing microcracks of intergranular type. (a) 2500x. (b) 4500x

stress, and the stress concentration at the crack tip accelerate hydrogen embrittlement and provide a driving force for crack propagation.

Figures 14 to 16 clearly show the view by SEM of the inter- granular decohesion resulting from cracking. The crack branching clearly indicates a constant crack growth rate over a range of applied stress intensities.

4. Conclusions

The severity and type of SCC depends on the sulfide con- centration in the seawater. The severity of SCC increases as the sulfide concentration increases. It was found that 90Cu- 10Ni is susceptible to sulfide cracking and hydrogen embrittlement. The hydrogen embrittlement is predominant at higher concen- trations of sulfide. The hydrogen attack is an indirect conse- quence of a series of electrochemical reactions in this sulfide-cupronickel alloy system. The stresses result in the rup- ture of the corrosion product film, thus enhancing sulfide stress cracking associated with the anodic dissolution in the low sul- fide concentration range and hydrogen embrittlement dominat- ing in the high sulfide concentration range. It was found that a synergism exists between sulfide and stress, which enhances the effect of the latter.

Acknowledgment The partial support of this work by the Kuwait University

Research Administration is gratefully acknowledged.

References 1. L. Schra and F.F. Groep, The ASCOR Test : A Simple Auto-

mated Method for Stress Corrosion Testing of Aluminum Al- loys, J. Testing Evaluation, Trans. of American Society for Testing and Materials, JTEVA, Vol 21 (No. 1), Jan 1993, p 44- 50

2. R.N. Parkins, Stress Corrosion Cracking--The Slow Strain Rate Technique, STP 665, ASTM, 1979, p 5-25

3. R.N. Parkins, Slow Strain Rate Testing--25 Years Experience, Slow Strain Rate Testing for the Evaluation of Environmentally In- duced Cracking: Research and Engineering Applications, ASTM STP 1210, R.D. Kane, Ed., ASTM, 1993, p 7-21

4. L.E. Eiselstein, B.C. Syrett, S.S. Wing, and R.D. Caligiuri, Cor- ros. Sci., Vol 23 (No. 3), 1993, p 223-239

5. A.M. Beccaria, G. Poggi, P. Traverso, and M. Giazza, Corros. Sci., Vol 32 (No. 11), 1991, p 1263-1275

6. J.N. Alhajji and M.R. Reda, Corros. Sci., Vol 34 (No. 1), 1993, p 163-177

7. J.N. Alhajji and M.R. Reda, Corros. Sci., Vo149, 1993, p 809 8. J.N. Alhajji and M.R. Reda, "The Effects of Some Common

Pollutants on the Corrosion of Copper-Nickel Alloys in Seawa- ter," The First UAB Symposium on Materials Science: Corro- sion and Corrosion Protection, United Arab Emirates, Dec 12- 14, 1993

9. M.R. Reda and J.N. Alhajji, "The Effect of Complexing Agents on the Corrosion of Copper/Nickel Alloys in Sulfide Polluted Seawater under Impingement Attack," 12th International Corro- sion Congress (Houston, TX), National Association of Corrosion Engineers, 1993

10. M.R. Reda and J.N. Alhajji, Br. Corros. J., Vo130 (No. 1), 1995, p 56-62

11. J.N. Alhajji and M.R. Reda, J. Electrochem. Soc., Vol 141 (No. 6), 1994, p 1432

12. D. Thompson, Material Research Study, Vol 1, 1961, p 108 13. J. Popplewell and T. Gearing, Corrosion, Vol 8, 1975, p 279-286 14. K. Habib and A. Husain, Stress Corrosion Cracking of Copper-

Nickel Alloys in Sulfide Polluted Natural Seawater at Moderate Temperature, Desalination, Vol 97, 1994, p 29-34

15. B.C. Syrett, Corros. Sci., Vol 21 (No. 3), 1981, p 187-209 16. C. Kato, H.W. Pickering, and J.E. Castle, J. Electrochem. Soc.,

Vol 3 (No. 6), 1984, p 1225 17. R. Francis, Br. Corr. J., Vol 20 (No. 4), 1985, p 175 18. C. Kato, B.G. Ateya, J.E. Castle, and H.W. Pickering, J. Electro-

chem. Soc., Vol 127, 1980, p 1890-1906 19. H.P. Hack and J.P. Gudas, Corrosion, Vol 35 (No. 6), 1979, p

259-265 20. J.P. Gudas and H.P. Hack, Corrosion, Vol 35 (No. 2), 1979, 67-

72 21. M.R. Reda and J.N. Alhajji, J. Univ. Kuwait (Sci.), Vo120, 1993,

p 171-186

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22. I.L. Rozenfeld, in Corrosion lnhibitors, (McGraw-Hill, 1981), p 276

23. C.D. Kim and B.E. Wilde, A Review of the Constant Strain Rate Stress Corrosion Cracking Test (Stress Cracking Corrosion-- The Slow Strain Rate Techniques), ASTM-STP 665, G.M. Ugian- sky and J.H. Payer, Ed., ASTM, 1979, p 97-112

24. B.C. Syrett, D.D. MacDonald, and S.S. Wing, Corrosion, Vol 35 (No. 9), 1979, p 409-422

25. J.H. Payer, W.E. Berry, and W.K. Boyd, Evaluation of Slow Strain-Rate Stress Corrosion Tests Results (Stress Cracking Corrosion--The Slow Strain Rate Techniques), ASTM-STP 665, G.M. Ugiansky and J.H. Payer, Ed., ASTM, 1979 p 61-77

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