Comparative evaluation of environment induced cracking of conventional and advanced steam turbine blade steels. Part 2: Corrosion fatigue

Corrosion Science 53 (2011) 503–512

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Comparative evaluation of environment induced cracking of conventional and advanced steam turbine blade steels. Part 2: Corrosion fatigue A. Turnbull ⇑, S. Zhou Materials Centre, National Physical Laboratory, Teddington, Middlesex TW11 0LW, UK

a r t i c l e

i n f o

Article history: Received 31 May 2010 Accepted 6 October 2010

Keywords: A. Steel A. Steam turbine C. Corrosion fatigue

a b s t r a c t Corrosion fatigue crack propagation rates have been determined for two steam turbine blade steels, PH13-8, a candidate steel for advanced turbines, and FV566, typical of conventional turbine blades. The testing was undertaken in simulated condensate environment, 300 ppb Cl and 300 ppb SO24 , at 90 °C using trapezoidal loading with a rise time of 20 min to simulate two-shifting (switching on- and off-load on a daily basis). Aerated and deaerated conditions were tested, the former being more representative of the retained aeration in the chamber as the system comes on load. In aerated solution at DK = 20 MPa m1/2 the cyclic crack growth rate for the PH13-8 steel was about a factor of two lower than that for the FV566 steel but the reverse was true for deaerated solution. In both cases, the difference in growth rates diminished with increasing DK. Although the cyclic crack growth rate is high the number of cycles per annum is small and the direct impact on overall life may be modest, but not insignificant. The observed cracking behaviour is best explained by a hydrogen assisted cracking mechanism. Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved.

1. Introduction

2. Experimental

Start-up and shut-down are the most potentially hazardous situations for plant operation because of the combined transient loading and environmental changes. In developing advanced turbines for converting steam to electricity it is not sufficient to consider simply on-load and off-load behaviour as a means of characterising likely service behaviour. The slow rising load associated with start-up together with residue of the off-load chemistry may impact on the initiation of cracks that would otherwise not be observed in laboratory testing with plain specimens under static load. Furthermore, repetition of start-up and shut-down (twoshifting) is increasingly desirable to respond to fluctuating power demand and to render older plant more economical. Thus, when introducing new blade steels into service, in new turbine systems or as retrofit, it is important to have a measure of the impact of such dynamic contributions to initiation and growth. Here, the focus is on simulating start-up and shut-down and assessing how the corrosion fatigue crack propagation rate for a proposed advanced steam turbine steel [1], PH13-8, compares with that for a conventional blade steel FV566. This study complements previous measurements of stress corrosion resistance of these two alloys [2].

The materials tested were a FV566 stainless steel and a PH13-8 stainless steel with composition as described in Table 1. The FV566 stainless steel had been annealed at 1050 °C for 1 h 45 min then air-cooled, tempered at 650 °C for 4 h and air-cooled. The specimens were stress relieved after manufacture at 600 for 2 h. The PH13-8 stainless steel had been solution annealed at 975 °C, aircooled to below 16 °C, tempered at 450 °C and air-cooled. As for the FV566 steel, the specimens were stress relieved in vacuum at 600 °C after manufacture. It was pertinent to determine the inclusion content of the steels tested to ensure that differences in cracking response could not be attributed to differences in inclusion content and its effect on local chemistry and crack advance. For the FV566 steel the inclusion density was determined to be 290 cm 2 with average inclusion size of 3.4 lm and for the PH13-8 steel the density was 280 cm 2 and the average size 2.8 lm, differences that would not be considered too significant. The microstructure of the FV566 steel was martensitic with a mean grain size of 27 ± 2 lm. For the PH13-8 steel the microstructure was also martensitic but with approximately 6.8% retained austenite (there was no detectable retained austenite in the FV566 steel). The grain size was 25 ± 1 lm. The mechanical properties were measured at ambient and at test temperature and are listed in Table 2. Compact tension (CT) specimens were made in accordance with ISO 7539-6 [3]. Side-grooves were not adopted. The thickness (B) and width (W) were 20 mm and 40 mm, respectively. The notch

⇑ Corresponding author. Tel.: +44 20 8943 7115; fax: +44 20 8614 0436. E-mail address: [email protected] (A. Turnbull).

0010-938X/$ - see front matter Crown Copyright Ó 2010 Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.10.001

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Table 1 Composition of steels (in mass %). Steel

C

Si

Mn

P

S

Cr

Mo

Ni

Cu

V

Al

Ti

Fe

FV566 PH13-8

0.11 0.05

0.23 0.15

0.71 0.05

<0.009 0.006

<0.003 0.002

11.69 13.2

1.73 2.24

2.71 8.14

/ 0.02

0.30 0.003

/ 1.02

/ 0.07

Bal Bal

r0.2/MPa

Material

T/°C

FV566

23 90

894 841

UTS/MPa 975 937

29 27

Elongation/%

PH13-8

23 90

1305 1263

1354 1287

19 18

depth, as measured from the loading line, was 14 mm. The specimens were fatigue pre-cracked to a depth of 6 mm from the root of the notch such that the total crack length, a, from the loading line (including the notch) was 20.0 mm (as measured at the surface in this case). In one case, FV566 steel in aerated 300 ppb Cl +300 ppb SO24 , a CT specimen with thickness of 15 mm and width of 30 mm was used. The notch depth was 7.5 mm and the total initial crack length was 13.5 mm. The range of K values tested was within the constraints of the plain strain requirements specified in ISO 7539-6 [3]. The pre-cracking was conducted according to ISO 7539-6 [3]. The pre-cracking fatigue frequency was 110 Hz and the stress ratio was 0.1. The maximum stress intensity at the final stage of the pre-cracking was 10.1 MPa m1/2. Comparative tests for the two alloys were conducted in aerated and deaerated 300 ppb Cl and 300 ppb SO24 (as sodium salts)1 to represent condensate chemistry under normal operating conditions [4], though there will be uncertainty as to the detailed chemistry (and degree of aeration) in the condensate during start-up. In some regions this will be more dilute but spraying of the last-row blades to limit heating due to ‘‘windage’’ effects could give rise to more concentrated solution, depending on the quality of the cooling water. Unfortunately, extended investigation of the effect of concentration on crack growth rates, similar to that reported for stress corrosion cracking [2], was not feasible. The test solution was prepared from analytical grade chemicals and high purity water with initial conductivity 0.06 lS cm 1 and was circulated to the test cell from a 22-litre reservoir (see Ref. [2] for full details of the loop and water chemistry control). The solution was pre-heated before entering the test cell. The temperature was maintained by heaters around each cell and controlled by thermocouples placed in each cell. Separate thermocouples were used to record the temperature near to each specimen. The test temperature was maintained at 90 ± 1 °C for direct comparison with previous work and because this is the temperature typical of early condensate formation on the low pressure turbine [4]. For tests in deaerated solutions, a very low oxygen level (typically below 2 ppb and always less than 5 ppb) was achieved by using stainless steel tubing, stainless steel cells, a stainless steel dosing pump and by continually passing nitrogen into the reservoir. Fresh solution was always pre-deaerated so that the low oxygen level was maintained during solution changes. For aerated solution, the oxygen concentration was maintained at 1.7–1.8 ppm by continually passing air into the reservoir at 90 °C. The volumetric flow rate through the system was 1.7 ml s 1 (about 6 l/h), corresponding to a linear flow rate of about 0.2 mm s 1 past the specimen. The corrosion potential in each cell was measured with respect to a saturated calomel reference electrode (SCE) to which all potentials quoted are referred. This electrode was held in a small 1

The concentrations in this paper relate to ppb (or ppm) by mass.

reservoir containing the test solution at ambient temperature and connected to the test cell via a tube also containing the test solution. A valve isolated the reservoir except when measurement was undertaken. Although the conductivity of the solutions is low (2.7 lS cm 1 for aerated 300 ppb Cl +300 ppb SO24 and 2.2 lS cm 1 for deaerated 300 ppb Cl +300 ppb SO24 ) the potential drop for the corrosion potential measurements was considered negligible, as a electrometer with high input impedance of 1  1014 X was used. A trapezoidal load cycle was employed, most commonly with a 20 min rise time (to simulate rapid start-up), 100 min hold time, 20 min unload time and 20 min at zero load. The choice of hold time is a balance between accelerating the testing and allowing any time dependent factors to exert an influence. Nevertheless, the effect of hold time on the crack growth rate was evaluated. The crack length was monitored using a multi-channel pulsed direct current potential drop (DCPD) system with high resolution and stability. The methodology was described previously [2,5] with the distinction that for these tests it was necessary to focus on measurements only at maximum load. In that context, 400 current cycles were applied each day over a time period of 75 min at maximum load and the data then averaged to obtain one potential drop value per day. At the end of the test the specimen was removed from the cell, washed and dried and then fatigued in air until final failure. The fracture surface was then examined after cleaning using ‘‘Super Clarke’s solution’’ (5 g/l of 1,3-Di-n-butyl-2-thiourea (DBT) in 18.8% HCl).

3. Results 3.1. FV566 steel 3.1.1. Corrosion potential measurements The time dependence of the corrosion potential of the FV566 steel in the corrosion fatigue tests is shown in Fig. 1. The corrosion potential was between 0.13 V(SCE) and 0.25 V(SCE) in aerated 300 ppb Cl and 300 ppb SO24 solution with a decreasing value

0.0 Aerated Deaerated

-0.1 Potential/V(SCE)

Table 2 Mechanical properties of the alloys tested.

-0.2 -0.3 -0.4 -0.5 -0.6 0

1000

2000

3000 Time/hours

4000

5000

Fig. 1. Corrosion potential of FV 566 blade steel in aerated and deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C.

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with exposure time suggesting an increased corrosion activity. In deaerated solution, the corrosion potential remained at about 0.59 V(SCE) throughout the test.

Fig. 3. Fracture surface of the FV566 specimen, highlighting the reasonably uniform crack front.

Holding time = 100 min da/dN = 1.3 μm/cycle

15.2 15.0 Crack length / mm

3.1.2. Crack growth rates The time-based crack growth rate in steam turbine steels is inherently small and thus it is most practical to determine the crack growth rate dependence on DK using a K-stepping method where the steady crack growth rate at a particular DK value is determined and the Kmax then stepped to a higher value and so on (minimum K maintained at zero). Fig. 2 shows the data obtained in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C. Stepping the load does seem to introduce some transient behaviour, almost as if the crack were temporarily blunted, before steady crack growth is attained. The cyclic crack growth rates are high, approaching 1.3 lm per cycle at the higher K values. To avoid the influence of temporary transients in growth rate in response to changing load or changing potential fitting to the raw data was made only when a steady growth rate was obtained. This principle was applied in all tests and the fitting error was always less than 5%. It was not possible to ratify the crack extension at each K value, as the crack front corresponding to the transition in DK value was not discernible. Nevertheless, the total crack extension measured on the fractured surface of 2.37 mm (average value at 9 locations) was in good agreement with that derived from the DCPD measurement, 2.41 mm (not all data shown in Fig. 2). The crack front was uniform, as shown in Fig. 3. Since the hold time in service is long and the use of 100 min an arbitrary means of accelerating the testing, it was pertinent to assess the impact of any hold time on the cyclic crack growth rate. The investigation of the effect of hold time on crack growth rates was made at a relatively high Kmax value of 40 MPa m1/2 to assess whether dynamic loading might influence the growth rate at maximum load just below or about the nominal KISCC (40–50 MPa m1/2). It is notable from Fig. 4 that almost the same cyclic growth rate was obtained with a hold time of 0 min as with 100 min but the growth rate was increased by about 50% when increasing the hold time to 300 min though with no further change at a hold time of 660 min. To assess whether this might reflect some evolution of the system with exposure time the hold time was returned to 100 min. The cyclic crack growth rate was as before. Switching to static, maximum, load for an extended period led to unusual transients before the growth rate reduced to a very low value of 0.14 mm/y. The reason for the apparent reduction in crack size in one case suggests

14.8

Static load Hold time = 300 min da/dN = 1.9 μm/cycle

Hold time = 660 min da/dN = 2.0 μm/cycle

14.6 Hold time = 0 min da/dN = 1.2 μm/cycle

14.4 Static load

Holding time = 100 min, da/dN = 1.3 μm/cycle

14.2 1000

2000

3000

4000

5000

6000

Time / hours Fig. 4. Effect of hold time on crack extension for FV 566 steel in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C. This test was a continuation of that shown in Fig. 2 with the initial value of DK = 40 MPa m1/2 and stress ratio zero.

14.3 1/2

ΔK = 40 MPa m

Crack length / mm

14.2

1/2

Δ K = 30 MPa m

14.1

14.0

da/dN = 1.3 μm/cycle 1/2 ΔK = 20 MPa m da/dN = 0.7 μm/cycle

13.9

13.8

da/dN = 0.4 μm/cycle

0

400

800

1200

1600

some shorting of the crack walls affecting the potential drop, perhaps associated with oxide bridging. It is evident from Fig. 4 that there is some time dependence of the crack length under static load even at long times so the crack had not quite arrested, at least within the 1000 h period. The results in deaerated solution are shown in Fig. 5 and suggest, by comparison with Fig. 2, no significant difference in the crack growth rate between aerated and deaerated solution. The effect of hold time on the cyclic growth rate of FV566 in deaerated 300 ppb Cl and 300 ppb SO24 was also evaluated. Increasing the hold time from 100 min to 300 min resulted in only a modest increase in the growth rate, from 1.4 lm/cycle to 1.6 lm/cycle.

Time / hours Fig. 2. Effect of DK on crack extension for FV 566 steel in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C. The initial value of DK was 20 MPa m1/2 and the stress ratio was zero.

3.1.3. Fractography The fracture surface of the FV566 steel formed in air at 90 °C is shown in Fig. 6. Scanning electron microscopy (SEM) examination

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22.0

Crack length / mm

21.5

thold = 300 min 1/2

Kmax = 40 MPa m

21.0

da/dN = 1.6 μm/cycle

1/2

Kmax = 30 MPa m da/dN = 1.4 μm/cycle

20.5

da/dN = 0.8 μm/cycle da/dN = 0.3 μm/cycle

20.0

0

1000

2000 3000 Time / hours

4000

Fig. 5. Effect of DK on crack extension for FV 566 steel in deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C. The initial value of DK was 20 MPa m1/2 and the stress ratio was zero.

of the fracture surface of the FV566 steel after exposure to the aerated solution proved challenging. The oxide on the fractured surface was not transparent to electrons and cleaning of the relatively thick oxide was necessary. This proved problematic, as

it was not possible to identify a solution that attacked the oxide without some damage to the substrate. The best option was ‘‘Super Clarke’s’’ solution (5 g/l of 1,3-Di-n-butyl-2-thiourea (DBT) in 18.8% HCl). However, comparison with a reference specimen showed that this created void-like defects after the extended exposure time required. The fracture surface appearance after cleaning is shown in Fig. 7 and highlights the observation of these voids. There was no evidence of the intergranular failure modes observed in the stress corrosion cracking tests [2] and no indication of quasicleavage modes. For the test in air at 90 °C, striations were observed on the fractured surface at relatively high stress intensity factor range (DK P 40 MPa m1/2), though only apparent at high magnification (5000). At DK  55 MPa m1/2 a crack growth rate could be estimated from the striation spacing (measured at different locations to ensure a reasonable average value) assuming that each striation corresponds to one cycle of fatigue loading. On this basis the growth rate was determined to be about 0.7 lm/cycle, which is in good agreement with the value of 0.6 lm/cycle derived from the DCPD measurements. For the test in aerated solution, the damaged fracture surface consequent upon cleaning precluded observation of striations. The specimen exposed to deaerated solution was much easier to clean as the film formed on the fractured surface was very thin and was removed in a time short compared to aerated solution.

Fig. 6. Fracture surface of FV566 steel in air at 90 °C; (a) DK = 20 MPa m1/2 and (b) DK = 60 MPa m1/2. The crack is growing from top-to-bottom.

Fig. 7. Fracture surface of FV566 steel fatigued in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C: (a) Kmax = 20 MPa m1/2 and (b) Kmax = 40 MPa m1/2. These fractogaphs relate to Fig. 2. The crack is growing from top-to-bottom. The apparent voiding is an artefact of the cleaning process and not intrinsic.

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Fig. 8. Fracture surface of FV566 steel fatigued in deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C: (a) pre-crack in air (PC) (Kmax = 10 MPa m1/2) to corrosion fatigue (CF) (Kmax = 20 MPa m1/2) and (b) CF (Kmax = 40 MPa m1/2) to fatigue in air (post crack) (Kmax = 50 MPa m1/2). These fractogaphs relate to Fig. 5. The crack is growing from top-tobottom.

Accordingly, no extraneous voiding was introduced. The cracking mode was transgranular, although there is some evidence of quasi-cleavage modes at low Kmax (Fig. 8a). It was difficult to discern striations confidently even at high magnification for both low and high DK and no quantitative measurement based on striation spacing was possible. 3.2. PH13-8 steel 3.2.1. Corrosion potential measurements For the PH13-8 steel, the corrosion potential in aerated solution (Fig. 9) was about 0.0 V(SCE), higher than that for FV566, reflecting the higher Cr and Ni content of the PH13-8 steel. When the solution was switched to the deaerated condition after 4104 h the corrosion potential decreased to 0.47 V(SCE) in 3 days. Two tests were conducted on the PH13-8 steel in sustained deaerated 300 ppb Cl and 300 ppb SO24 solution. In Test 1 of Fig. 9 the initial corrosion potential was 0.57 V(SCE) but gradually increased to about 0.32 V(SCE). At 3100 h, the potential decreased to reach a quasi-stable value of 0.46 V(SCE). In Test 2, the initial corrosion potential was 0.49 V(SCE) and gradually increased to about 0.28 V(SCE). At 4200 h, the solution was replaced by aerated solution, resulting in an increase in corrosion potential, gradually

22.0

Aerated Deaerated, Test 1 Deaerated, Test 2

Deaerated

0.0

1/2

Crack length / mm

Potential/V (SCE)

0.2

reaching a stable value of 0.06 V(SCE), similar to that measured in the constant aerated solution test. The drift in corrosion potential to more positive values for the PH13-8 steel in deaerated solution was surprising since the oxygen concentration in the tests with deaerated solution remained below 5 ppb. It had been noted in a test with FV566 steel that while the corrosion potential of this steel in deaerated solution was sustained at about 0.59 V(SCE) the potential of the 316L stainless steel containment vessel measured at the same time was about 0.15 V(SCE); i.e. a relative high potential can be achieved even in deaerated solution depending on the alloy composition and solution conductivity. This observation gives some measure of confidence that the drift in potential to more positive values for the PH13-8 steel was not spurious and is related to the higher alloying content of this steel combined with the low anion content of the solution. Nevertheless, it does imply a reduced passive current density or some increased activity for reduction of water with time, or both. The subsequent decrease in potential of the PH138 steel is more difficult to explain. It suggests an increase in corrosion activity perhaps reflecting a decreased potential in the crack though there was no particular change in the loading conditions at that time period. It seems less likely that this is due to crevice attack in the contact area at the loading pins as there is no differential aeration cell to drive the activity.

Deaerated

-0.2 Aerated

-0.4

ΔK = 42 MPa m

21.5

1/2

ΔK = 32 MPa m

da/dN = 1.9 μm/cycle

1/2

ΔK = 21 MPa m

da/dN = 1.0 μm/cycle

21.0

da/dN = 0.6 μm/cycle da/dN = 0.2 μm/cycle

-0.6

da/dN = 0.1 μm/cycle

0

2000

4000 Time / hours

6000

Fig. 9. Corrosion potential of PH13-8 steel in aerated and deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C. At about 4000 h the initially aerated test solution was switched to a deaerated test solution and in the case of Test 1 the initially deaerated solution was replaced with aerated solution.

20.5

0

1000

2000 3000 Time / hours

4000

5000

Fig. 10. Effect of DK on crack extension for PH13-8 steel in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C. The initial value of DK was 16 MPa m1/2 and the stress ratio was zero.

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3.2.2. Crack growth rates The crack growth rate for the PH13-8 steel in aerated 300 ppb Cl and 300 ppb SO24 as a function of DK is shown in Fig. 10. There is some small fluctuation of the crack growth rate for lower DK values but the values were reasonably steady. The most interesting observation is the clear increase in crack growth rate in converting from aerated solution to deaerated solution. The crack growth measurements obtained in a separate test in sustained deaerated solution are shown in Fig. 11. Notably, the growth rate is unsteady and decreasing for at least 500 h at a Kmax of 31.4 MPa m1/2, which coincided with a drift in corrosion potential to more positive values (Fig. 9). Adopting a hold time of 300 min appeared to result in a discernible, though modest, increase in cyclic crack growth rate but, since the DK was slightly greater, too much emphasis on this observation is not merited. A further test was conducted in deaerated 300 ppb Cl and 300 ppb SO24 solution at 90 °C with a initial Kmax of 20 MPa m1/2 and hold time of 0 min (Fig. 12). At 901 h, the hold time was increased to 100 min and there was no change in the crack growth rate. The crack growth rate at 20 MPa m1/2 was half of that in the previous test at the almost same K value (Fig. 11), probably because the corrosion potential, > 0.30 V(SCE), of the steel in this test (Test 2 in Fig. 9) was higher than that in the previous test,

1/2

t hold = 300 min, Kmax = 44.0 Mpa m

22 Crack length / mm

1/2

K Max = 42.0 MPa m

da/dN = 2.0 μm/cycle

1/2

KMax = 31.4 MPa m

21

da/dN = 1.8 μm/cycle da/dN = 0.8 μm/cycle da/dN = 0.6 μm/cycle

20

0

1000

2000

3000

4000

Time / hours Fig. 11. Effect of DK on crack extension for PH13-8 steel in deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C. The initial value of DK was 21 MPa m1/2 and the stress ratio was zero. This relates to Test 1 of Fig. 9.

22.5

thold = 300 min

O2 = 1.7 ppm

Crack length / mm

22.0

which was more negative than 0.50 V(SCE) (Test 1 in Fig. 9). The crack growth rate at 30 MPa m1/2 and 40 MPa m1/2 is also considerably lower than that in the previous test (Fig. 11), possibly reflecting a consistently higher corrosion potential in the second test. It is also interesting to note that there was no change in the crack growth rate when the solution was aerated, although the corrosion potential increased from 0.28 V(SCE) to 0.06 V(SCE), suggesting that the limit of crack-tip polarisation had already been achieved at the more negative potential. When the hold time was changed from 100 min to 300 min (aerated solution) there is an inherent change in the time-based growth rate, as seen in Fig. 12 for example, but there was no change in the cyclic crack growth rate. 3.2.3. Fractography The oxide film on the fracture surface of the PH13-8 steel was sufficiently transparent to electrons to obviate the need for cleaning. The fracture surface formed in air at 90 °C is shown in Fig. 13 and that formed during exposure to aerated 300 ppb Cl and 300 ppb SO24 is shown in Fig. 14. Striations were clearly visible on the fracture surface in the tests conducted in aerated solution at DK = 42 MPa m1/2 even at low magnification though less obvious at DK = 16 MPa m1/2. Striations were also apparent on the fractured surface formed in air at 90 °C though readily apparent only at high magnification (5000) in the region at DK = 55 MPa m1/2. The crack growth rate measured from the striation spacing was about 0.7 lm/cycle, almost the same as the value derived from the DCPD measurements. The growth rate measured from the striation spacing on the fracture surface formed in aerated solution at DK = 40 MPa m1/2 was in the range 1.3 lm/cycle to 1.4 lm/cycle (measured at several locations), consistently higher than the average value of 1.0 lm/cycle derived from the DCPD measurements. This observation suggest that not every cycle of loading produces crack extension along the whole crack front. This may be associated with pinning of the crack locally, by residual ligaments perhaps. There was also an indication that the fracture was more brittle in the environment at the lower DK value. The fracture surface obtained in deaerated solution is shown in Fig. 15. Striations can be observed at DK = 42 MPa m1/2 but not as marked as in aerated solution. However, the most notable feature is the very brittle zone at DK = 20 MPa m1/2. While it could represent a sensitivity of the fracture mode to the stress intensity factor, the possibility is also that this is some inhomogeneity in the material microstructure. To resolve this, the steel was sectioned parallel to the fracture surface to provide a slice of about 2 mm thick. Microhardness measurements were then made as a function of distance from the pre-crack with the Scanning Indenting Microhardness Measurement [6] tool using a 0.5 N load and 20 lm spacing of indents. No difference in microhardness was apparent.

1/2

KMax = 40 MPa m 1/2

21.5

KMax = 30 MPa m thold = 100 min

da/dN = 1.1 μm/cycle

4. Discussion da/dN = 1.1 μm/cycle da/dN = 1.2 μm/cycle

da/dN = 0.6 μm/cycle

21.0

da/dN = 0.3 μm/cycle da/dN = 0.3 μm/cycle, (thold = 0 min)

20.5

0

1000

2000

3000 4000 Time / hours

5000

6000

Fig. 12. Effect of DK on crack extension for PH13-8 steel in deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C. The initial value of DK was 21 MPa m1/2 and the stress ratio was zero. This relates to Test 2 of Fig. 9.

Perhaps the most striking feature of the corrosion fatigue results, by comparison with the stress corrosion data, is the significant crack growth rate in deaerated solution, in the case of PH13-8 steel being higher than that in aerated solution (Fig. 10). The time-dependence of the crack growth for the PH13-8 steel in deaerated solution at the same DK in Fig. 11 reflects the response of crack growth rate to the drift in corrosion potential from 0.5 V(SCE) soon after immersion to just above 0.3 V(SCE) at longer exposure times. The maximum difference in cyclic growth rate for this steel at the same DK was a factor of 2 with the higher growth rates being associated with the more negative potential of 0.5 V(SCE). However, the sensitivity of growth rate to increase in

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Fig. 13. Fracture surface of PH13-8 steel in air at 90 °C: (a) DK = 20 MPa m1/2 and (b) DK = 40 MPa m1/2. The scale bar is 30 lm and the crack is growing from top-to-bottom.

Fig. 14. Fracture surface of PH13-8 steel in aerated 300 ppb Cl and 300 ppb SO24 at 90 °C: (a) transition between pre-cracked region at DK = 10 MPa m1/2 to growth in the environment at an initial DK of 16 MPa m1/2; (b) growth at DK = 42 MPa m1/2.

Fig. 15. Fracture surface of PH13-8 steel in deaerated 300 ppb Cl and 300 ppb SO24 at 90 °C: (a) DK = 20 MPa m1/2 (b) growth at DK = 40 MPa m1/2.

potential tends to diminish as the potential drifts more and more positive suggesting attainment of a limiting crack-tip potential. The slow crack growth rate in these corrosion fatigue tests limits the extent of repeatability checks. Nevertheless, comparison of Figs. 10 and 11 for deaerated solution at DK  40 MPa m1/2 suggests consistent measurement of crack growth rate in separate tests on PH13-8 steel. Repeatability of the crack growth rate on the same specimen of FV566 steel after changing and switching

back the test conditions (Fig. 4) also is reassuring. Thus, it may be deduced that the trend in data with respect to the effect of aeration is real. The high crack growth rate with cyclic loading in deaerated solution compared to the almost insignificant growth rate under static load conditions [2] can be rationalised on the basis of the mechanical contribution to film rupture at the crack tip insofar as there is then less reliance on attaining an aggressive crack-tip

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chemistry and a critical balance of crack tip creep rate and refilming kinetics. For deaerated low conductivity solution the crack-tip potential will be relatively low and the crack-tip solution will correspondingly tend to be alkaline as there is no imbalance of the anodic and cathodic reactions in the crack as there would be in aerated solution. The mechanically driven change in crack-tip opening displacement (CTOD) exposing bare metal combined with the kinetics for water reduction at the crack tip will be sufficient to induce hydrogen uptake and accelerate the crack growth. An explanation based on crack-tip dissolution kinetics is untenable, especially when considering the enhanced local potential drop that will almost certainly be induced by exposure of fresh metal at the crack tip. Thus, there is a clear implication from the tests in deaerated solution that the mechanism of failure is associated with hydrogen. However, an explanation for the different response of the crack growth rates for the two alloys to the extent of aeration of the solution is required; i.e. the insensitivity for the FV566 steel but some dependence for the PH13-8 steel. For these dilute low conductivity solutions, it would be expected that crack-tip polarisation under fatigue loading conditions would be very limited, though by implication some does occur for the PH13-8 steel as evidenced by the reduced crack growth rate with increase in potential of the PH13-8 steel shown in Fig. 11. As noted also, there would be some transient partial decoupling of the tip potential from the external polarisation in response to the significant disruption of the cracktip film during rising load as the potential on the crack walls becomes preferentially more negative to balance the crack-tip anodic current. Such decoupling was indicated in crack-tip potential measurements for a low alloy turbine disc steel under these conditions [7]. This would tend to encourage hydrogen entry. A simple explanation for the different sensitivities of the two alloys to testing in aerated and deaerated solution, and hence corrosion potential, could be that the higher strength PH13-8 steel is just more sensitive to differences in extent of hydrogen uptake and thus will be more responsive than the FV566 steel to the small changes in crack-tip potential induced by changes in the corrosion potential. While plausible, a more substantive rationalisation for the effect of potential may relate to the combination of a smaller value of DCTOD for the higher strength steel, i.e. less new surface

each cycle, the faster refilming kinetics and the lower passive current density on the crack walls expected for the PH13-8 steel. The net effect would be less total current flowing from the crack and more likelihood in these dilute solutions of crack-tip polarisation. Thus, an effect of corrosion potential is observed for the high strength, higher-alloyed steel but not for the low strength steel, for which there is projected to be no or very limited crack-tip polarisation. Nevertheless, a limiting crack-tip potential will be obtained for the PH13-8 steel. Thus, in Fig. 12 it can be observed that there is no difference in crack growth rate upon aerating the previously deaerated solution because the corrosion potential in the former after over 4000 h exposure was already quite positive at about 0.28 V(SCE). At this high potential, no further polarisation of the crack tip would be possible. Implicit also in this evaluation is a reduced amount of hydrogen generation per cycle in the case of the high strength PH13-8 steel. The crack growth rate is actually slightly greater at similar corrosion potentials to that for the FV566 steel (Fig. 16) but the difference was small compared with the difference in stress corrosion cracking crack growth rates. Rationalising the effect of strength level on corrosion fatigue of these two alloys is complex. While it is inferred that there is less hydrogen generated per cycle for the PH13-8 steel, the hydrogen generated may be more effectively absorbed and retained at the crack tip as a consequence of the greater hydrostatic stress component for the higher strength steel. The two alloys may also exhibit different degrees of material response to cyclic loading in terms of their tendency for cyclic softening or hardening and the dislocation cell structures formed may be different. The retained austenite in the PH13-8 SS will also likely undergo phase transition to martensite. Furthermore, in this region of growth rates there is no explicit stress corrosion component and ‘‘true’’ corrosion fatigue crack growth rates may not scale on strength level in the manner expected for static loading [8]. A curious feature of these measurements has been the effect of hold time. Hold time effects can be associated with the time evolution of crack-tip chemistry and potential in response to dynamic load and/or hydrogen ingress and transport at the crack-tip. The hold time of 100 min adopted in most tests was based on two observations. The crack-tip potential of a turbine disc steel in deaerated water at 90 °C reached a stable value within 100 min

-6

10

-7

da/dN / m/cycle

10

-8

10

-9

10

-10

10

FV566: Air Aerated solution, Ecorr ~ -0.19 V (SCE) Deaerated solution, Ecorr ~ -0.59 V (SCE)

PH13-8: Air Aerated solution, Ecorr ~ -0.00 V (SCE) Deaerated solution Ecorr ~ -0.49 V (SCE)) Ecorr ~ -0.47 V (SCE) Ecorr ~ -0.36 V (SCE) Ecorr ~ -0.31V (SCE) Ecorr ~ -0.27V (SCE)

10

100 1/2

ΔK / MPa m Fig. 16. Comparison of crack growth rates of FV566 and PH13-8 blade steels in 300 ppb Cl and 300 ppb SO24 at 90 °C under trapezoidal loading conditions (hold time 100 min; R = 0) and in air at high frequency. For tests in air (at 90 °C), the growth rate in the Paris law region could be expressed by 9.4  10 12 DK2.7 m/cycle for FV566 steel and 2.2  10 11 DK2.6 m/cycle for PH13-8 steel.

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after increasing the load from zero to maximum in a fatigue cycle [7]. Furthermore, the results of Endo et al. [9] also indicated limited influence of hold times at times greater than 1 h to 2 h. The investigation of the effect of hold time on crack growth rates in deaerated solution was limited. Nevertheless, no effect was discernible for PH13-8 steel (Fig. 12) when the hold time was increased from 0 min to 100 min and the increase in growth rate when the hold time was increased from 100 min to 300 min is sufficiently small (about 11%) to be of the order of uncertainty in the fitting to the raw data. Similarly, there was only a very modest effect on growth rate for the FV566 steel in deaerated solution. There was no effect of hold time on the cyclic crack growth rate of the PH13-8 steel in aerated solution, at least in extending the hold period from 100 min to 300 min, but the lack of reference data at zero hold time limits deduction. However, the data of Fig. 4 in relation to FV566 steel in aerated solution are challenging to explain both with respect to hold time effects and also the transition to static load for a very extended period. For the former, the problem is to explain an absence of hold time effects between 0 min and 100 min but an increase in cyclic crack growth rate when the hold time was increased from 100 min to 300 min. As noted, changes in corrosion potential are not a factor and the data appear robust. However, it is difficult to conceive of any process, such as recovery of crack-tip chemistry that is ineffective during the first 100 min but begins to have a role later, especially as there did not seem to be any effect of aeration on growth rate at 100 min for this steel. For the moment it remains unexplained. More concerning perhaps is the inconsistent behaviour on transforming to static load and then leaving it under these conditions. Unusually long transients in growth rates are observed that are likewise difficult to explain and perhaps question the basis of the simple accelerated test using 100 min hold time, though it is difficult to conceive of a more productive approach for laboratory testing. Perkins and Bache [10] tested plain FV566 steel specimens in aerated 1 ppm chloride at 120 °C and noted that a dwell period at maximum load appeared to increase the number of cycles to failure. Indeed they found the least damaging waveform was a slow ramp and a dwell. However, it is possible that these factors relate to the development of pitting insofar as pitting for this system, at this low chloride concentration, is driven by the fatigue process; i.e. there is no pitting in the absence of mechanical loading. A greater rate of mechanical damage may be conducive to the more rapid development of pits, though the subsequent cyclic crack growth might respond in an inverse way. In contrast to stress corrosion cracking, the failure mode in this range of DK was wholly transgranular for both aerated and deaerated solutions. There was no indication of intergranular cracking at the Kmax of 40 MPa m1/2, the limit of testing, reflecting perhaps the relatively high KISCC threshold for both steels in the very dilute solution tested and suggesting that this threshold is not significantly reduced in response to fatigue loading. Intergranular cracking modes have been observed in corrosion fatigue of blade steels at lower DK values [11] with a peak at about 18 MPa m1/2. However, the frequency of testing was much higher (60 Hz) and the sodium chloride concentration was 3% so direct comparison is not informative. Nevertheless, the possibility of intergranular modes at lower DK values than we have measured cannot be discounted though it is surmised that such observation may be linked in some way to a lower threshold for static crack modes in the 3% NaCl solution, noting the sensitivity of percentage intergranular factor to NaCl concentration. However, in their study of FV566 steel (plain specimens) Perkins and Bache [10] also noted mixed mode cracking at lower concentration of chloride. In aerated 1 ppm chloride they observed mixed transgranular/intergranular modes and then transgranular cracking prior to tensile overload. Interestingly, they noted that in some tests with ‘‘zero’’ oxygen and no chloride there was evi-

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dence of mixed quasi cleavage crack propagation. Such quasi-cleavage modes were not apparent in our tests for this steel. The fractography of the PH13-8 steel at DK 20 MPa m1/2 suggests quite brittle behaviour in deaerated solution but this is much less apparent in aerated solution and in both environments at DK = 40 MPa m1/2 brittle facture modes are not apparent. The striking nature of the brittle fracture in deaerated solution gave rise to the possibility of a local hard zone in the material. However, microhardness measurements indicated no transition in microstructure associated with this brittle zone, leaving an explanation for the discreteness of this observation difficult to propose. The corrosion fatigue literature for steam turbine steels was reviewed previously by Zhou [12]. Much of the crack propagation data was obtained at a higher frequency than adopted for our study and often in quite concentrated environments (e.g. 3.5% NaCl or higher), so direct comparison of crack growth rates is not so meaningful. Nevertheless, the results of Cho et al. [13] with frequency 0.5 Hz in 3.5% NaCl and in distilled water suggest crack growth rates between 10 7 m/cycle and 10 6 m/cycle, of the same order as measured in this study. 4.1. Engineering implications In relation to simulating repetitive start-up and shut-down in service, the behaviour in aerated solutions is perhaps most relevant on the basis that most damage results from the loading part of the cycle and oxygen depletion will be slow coming on load. As such, the advanced steel, PH13-8, performs better than FV566 with respect to resistance to corrosion fatigue crack propagation in normal water chemistry conditions though the decrease in crack growth rate is less than or about 50%. This is in contrast to the stress corrosion cracking response in which the growth rate was about a factor of 10 higher for the advanced steel. For this zero stress ratio a threshold DK similar to that in air might conservatively be assumed with crack closure dominating. There has been insufficient study to determine if crack growth is sustained at maximum load for Kmax values lower than the KISCC measured in static load testing but the limited evidence suggests that this is not significant. For both alloys the cyclic crack growth rates are high but the number of cycles per year will be modest. Assuming two-shifting 5 days per week and shut down over the weekend, the annual number of cycles will be about 260. For the PH13-8 steel with a growth rate in aerated solution of about 2.0 lm/cycle at DK of about 40 MPa m1/2 (allowing for hold time effect) this would represent an annual crack extension of about 0.5 mm. It is perhaps unlikely that new plant operating for 40 years would do so under two-shifting conditions. However, two-shifting might provide the basis for generating the precursor to stress corrosion cracking, though the latter would be significant only for aerated conditions and thus growth would be limited to relatively short transient periods. 5. Conclusions (1) The fatigue crack growth rates in air were very similar for the PH13-8 and the FV566 steels. (2) In simulated condensate solution at 90 °C the relative crack growth rates of the different steels varied for deaerated and aerated solution but in both cases exceeded the air data by a factor of about 5 or more, depending on the DK value. (3) For aerated solution the cyclic crack growth rate for the PH13-8 steel was about a factor of two lower than that for FV566 steel at 20 MPa m1/2 but the reverse was true for deaerated solution and the relative difference diminished with increase in DK.

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(4) There was no difference in crack growth rates for FV566 steel under aerated or deaerated conditions. This can be explained by a lack of crack-tip polarisation in the dilute solution tested, a consequence of the potential drop induced by the local high current density associated with the new crack surface produced in each loading cycle. (5) The contrasting effect of aeration level on crack growth rates for the high strength PH13-8 steel, compared with FV566 is considered to be due to crack-tip polarisation being more readily possible since the cyclic crack-tip opening displacement associated with the higher strength steel is smaller (less new surface area at the crack tip per cycle), the refilming kinetics are likely to be faster, and the passive film current density less for this higher-alloyed steel. (6) The observation of high crack growth rates in deaerated solution relative to those obtained in air suggests that hydrogen assisted cracking is the prevalent failure mechanism. The contrast with the stress corrosion behaviour for which cracks arrested in deaerated solution can be attributed to the dominant influence of the mechanical loading in generating bare, reactive, surface at the crack tip. (7) An unusual effect of hold time on crack growth rates for the FV566 steel in aerated solution was observed that could not be rationalised. (8) Both steels exhibited transgranular cracking for the DK values tested. For the PH13-8 steel, distinct cleavage facets were observed at a DK of about 20 MPa m1/2 but were not apparent at higher DK. No indication of any unusual microstructure was detected that might account for the discreet nature of it.

Acknowledgements This work was conducted as part of a joint venture between the United Kingdom Department of Industry, Universities and Skills and an industrial Group comprising S. Fenton (E. On UK), D. Gass (Siemens), P. McIntyre (IOMMM), N. Shaw (RWE npower), M. Tookey (British Energy), and S. Osgerby (Alstom Power). References [1] T. Furuse, R. Masumoto, Y. Yamamoto, Y. Kadoya, H. Ooyama, Development of high strength ph-stainless steel for steam turbine forged long blade, in: Proceedings 17th International Forgemasters Meeting, Santander, Spain, 2008, pp. 466–467. [2] A. Turnbull, S. Zhou, Comparative evaluation of environment induced cracking of conventional and advanced steam turbine blade steels. Part 1: stress corrosion cracking, Corros. Sci. 52 (2010) 2936–2944. [3] ISO 11782-2:1998, Corrosion of Metal and Alloys – Corrosion Fatigue Testing, Part 2: Crack Propagation Testing Using Precracked Specimens, International Standards Organisation, Geneva, 1998. [4] S. Zhou, A. Turnbull, Corros. Eng. Sci. Tech. 38 (2003) 97–111. [5] S. Zhou, A. Turnbull, Corrosion 62 (2006) 508–513. [6] J.D. Lord, B. Roebuck, R. Morrell, T. Lube, Mater. Sci. Tech. 26 (2010) 127–148. [7] A. Turnbull, G. Hinds, S. Zhou, Corros. Sci. 46 (2004) 193–211. [8] R.P. Gangloff, Private Communication, April 2010. [9] T. Endo, H. Itoh, Y. Kondo, H. Karato, Material aspects for the prevention of environmentally assisted cracking in low pressure steam turbine, in: PWR, The Steam Turbine Generator Today: Materials, Flow Path Design, Repair and Refurbishment, vol. 21, ASME International, New York, 1993, pp 75–82. [10] K.M. Perkins, M.R. Bache, Int. J. Fatigue 27 (2005) 1499–1508. [11] R. Ebara, T. Yamada, H. Kawana, ISIJ Int. 30 (1990) 535–539. [12] S. Zhou, Environment assisted cracking of turbine blade steel – a review, NPL Report DEPC-MPE 033, 2007. [13] S.Y. Cho, C.H. Kim, D.H. Bae, Key Eng. Mater. 993 (2000) 183–187.