Enhanced corrosion resistance of ultrafine-grained AZ61 alloy containing very fine particles of Mg17Al12 phase

Corrosion Science 75 (2013) 228–238

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Enhanced corrosion resistance of ultrafine-grained AZ61 alloy containing very fine particles of Mg17Al12 phase H.S. Kim a, W.J. Kim b,⇑ a Department of Nanomaterials Engineering, College of Nanoscience & Nanotechnology, Pusan National University, 50 Cheonghak-ri, Samnangjin-eup, Miryang-si, Kyongnam 627-706, Republic of Korea b Department of Materials Science and Engineering, Hongik University, Mapo-gu, Sangsu-dong 72-1, Seoul 121-791, Republic of Korea

a r t i c l e

i n f o

Article history: Received 24 January 2013 Accepted 31 May 2013 Available online 12 June 2013 Keywords: A. Magnesium B. Weight loss B. Polarization B. TEM C. Passive films

a b s t r a c t Ultrafine-grained (1.1–1.5 lm) AZ61 alloys containing nanoscaled b-Mg17Al12 phase particles (70– 140 nm) were prepared using high-ratio differential speed rolling (HRDSR) and their corrosion behaviours were studied in a 0.1 M NaCl solution. The grain size reduction by HRDSR improved the corrosion resistance by enhancing passivity of the surface film. Post-annealing further increased the corrosion resistance by decreasing dislocation density in matrix. When significant grain growth took place, however, the corrosion resistance was decreased because of the pronounced negative effect of the increase in grain size. Refinement of b phase to nanoscale size decreased the susceptibility to microgalvanic corrosion. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction Mg and Mg alloys are light metals that have great potential for application in electronics, transportation, and biomedical fields, wherein light weight and biocompatibility are critical issues. However, these materials have low strength and poor corrosion resistance, which limit their applications. Grain refinement is an important method for enhancing the strength of metals. It is particularly effective for enhancing the strength of Mg and Mg alloys because of their high Hall–Petch strengthening coefficient [1]. Equalchannel angular pressing (ECAP), which is a representative method of severe plastic deformation (SPD), has been popularly used to effectively decrease the grain size in Mg and Mg alloys [2–7]. ECAP is, however, a discontinuous process and is inadequate for handling materials in sheet form. Kim et al. [8–10] have shown that high-ratio differential speed rolling (HRDSR), in which a large amount of shear strain is induced during rolling, is highly efficient for refining grains in magnesium alloys in sheet form. In addition to chemical factors such as chemical composition, structural factors such as the amount and distribution of secondary phases [11,12], texture [13,14], dislocation density, and grain sizes [15–19] have been shown to significantly affect the microgalvanic corrosion and corrosion barrier properties. Kim et al. [8,9] recently used HRDSR to develop very-finegrained microstructures in AZ61 (Mg–6Al–1Zn) alloys containing ⇑ Corresponding author. Tel.: +82 2 320 1468; fax: +82 2 325 6116. E-mail address: [email protected] (W.J. Kim). 0010-938X/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2013.05.032

nanoscaled particles of b-Mg17Al12 phase. The HRDSR-processed AZ61 alloys showed very high yield strength over 380 MPa primarily as a result of grain refinement [9]. The materials also exhibited excellent superplasticity at elevated temperatures [8] because of the very fine b particles at the grain boundaries of the matrix, which enhance the thermal stability of the microstructure by hindering grain boundary migration during grain growth. The purpose of the present work is to examine the corrosion behaviour of the HRDSR-processed AZ61 alloys and to understand the structural factors affecting the corrosion properties in a NaCl environment. According to the results, the corrosion resistance of the AZ61 alloy can be greatly improved by reducing the size of grains, the density of dislocations, and the size of b-Mg17Al12 particles. 2. Experimental procedures 2.1. Preparation of materials A 50 mm wide  1.6 mm thick AZ61 plate was used as the starting material in this study. Its chemical composition (wt.%) was 6.43Al, 1.09Zn, 0.01Mn, 0.01Si, 0.001Cu, 0.004Fe, 0.001Ni and balance Mg. The plate was rolled using a rolling mill with rolls 300 mm in diameter. The ratio of the speed of the top and bottom rolls was maintained at either 2 or 3; the top roll rotated counterclockwise at either 6 or 9 rpm, and the bottom roll rotated clockwise at 3 rpm. A cold AZ61 plate was fed through the hot rolls at 150 °C to decrease the total thickness of the plate by 66% or 67% in two steps. The rolls were maintained at 150 °C, and the plate

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(a)

(b)

50 μm

5 μm

Fig. 1. (a) SEM micrograph of as-received AZ61. (b) SEM micrograph showing divorced eutectic b phase at grain boundary.

(a)

(b)

50 μm

50 μm

(c)

(d)

1 μm

50 μm

Fig. 2. Optical micrographs of (a) as-received, (b) HRDSR 2SP, and (c) HRDSR 3SP samples, (d) SEM micrograph of HRDSR 3SP sample.

was initially rolled to a thickness of 1.1 mm and then to a final thickness of either 0.55 or 0.53 mm. The plate was re-rolled without rotating it, i.e., the direction of rolling used during the second rolling was the same as that used during the first one. The sheets obtained after the second rolling at the speed ratios of 2 and 3 will be hereafter referred to as HRDSR 2SP and HRDSR 3SP, respectively. Samples of HRDSR 2SP were annealed at 240, 280, 320, or 400 °C for 2 h to examine the effect of annealing on the corrosion properties of the samples. 2.2. Examination of microstructure An optical microscope (OM) and a scanning electron microscope (SEM, JSM-5600LV) were used to observe the microstructures in the transverse direction (TD)–rolling direction (RD) or RD-normal direction (ND) sections of the samples after etching by a solution composed of 5% picric acid, 13% acetic acid, 12% distilled water and 70% ethanol. To define the thickness layer of the sheet, a parameter s = ± 10Dx/(do/2) was introduced, in which Dx represents the distance between the respective layer and the sample

centre and do/2 is the half-thickness (i.e., s = 0 for the centre layer, s = 10 for the upper surface layer in contact with the high-speed roller, and s = 10 for the lower surface layer in contact with the low-speed roller). A field-emission transmission electron microscope (FE-TEM, JEM 2001F/200 kV) equipped with an energy-dispersive X-ray spectroscope (EDS) was used to examine the detailed microstructures on the TD–RD sections (at s = 0) of the samples. Solutions composed of 60% methanol, 30% glycerine, and 10% nitric acid were used for jet polishing, and a BAL-TEC RES 101 was used for ion milling. Focused ion beam (FIB) milling (Helios 200 NanoLab FIB) was used to prepare the RD–ND sections of the samples for high-resolution TEM observation (Tecnai F20 TEM/200 kV). Cu Ka radiation was then used to measure X-ray pole figures on the surfaces (at s = 9) of the samples in order to characterize the texture of the samples. 2.3. Electrochemical measurements Potentiodynamic polarization experiments were performed using a potentiostat/galvanostat (GAMRY series, type G300) in

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(a)

(b)

500 nm

(c)

500 nm

(d)

500 nm

500 nm

(f)

(e)

500 nm

500 nm

Fig. 3. TEM micrographs of (a) the as-received AZ61 alloy and (b) HRDSR 3SP. (c) Bright- and (d) dark-field TEM micrographs of HRDSR 2SP. (e) Bright- and (f) dark-field TEM micrographs of HRDSR 3SP.

Fig. 4. Element maps for Al and Mn overlapped on a TEM image of HRDSR 3SP. Green and yellow patch indicate X-ray signals from Al and Mn, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

solutions at room temperature (23 ± 1 °C). Samples were ground up to s = 9 with emery paper (4000 grit), followed by a polishing

treatment with a diamond suspension. After polishing, the samples were ultrasonically cleaned in ethanol. The electrochemical cell was equipped with an Hg–Hg2Cl2–saturated KCl electrode, a graphite counter electrode, and a Luggin capillary. The cleaned samples were moulded in TeflonÒ with a 10-mm-diameter window and then exposed to 0.1 M NaCl solutions. Potentiodynamic polarization tests were performed at a scan rate of 5 mV/s in the range 1.9 to 0.7 VSCE. A lower scan rate of 1 mV/s was initially used, but it provided unstable measurements in obtaining anodic branches of the polarization curves from some samples. The samples were incubated for 60 min to stabilize the initial conditions before starting the potential scans. From the obtained polarization curves, the cathodic Tafel slope (bc) was calculated by a linear fitting between 1.8 and 1.6 VSCE and the corrosion current density (icorr) was measured by the extrapolation of cathodic Tafel region to the corrosion potential (Ecorr).

2.4. Weight loss measurements The rate of corrosion was calculated from the weight loss caused by corrosion. Before the immersion test, the surfaces of the samples were ground using emery papers (up to 4000 grit),

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(a)

(b)

20 μm

(c)

20 μm

(d)

20 μm

20 μm

Fig. 5. Optical photographs of HRDSR 2SP samples annealed at (a) 240, (b) 280, (c) 320, or (d) 400 °C for 2 h.

followed by a polishing treatment with a diamond suspension. The finally prepared samples of the as-received AZ61, HRDSR 2SP, and HRDSR 3SP had the dimensions of 10 (length)  10 (width)  1.0 (thickness) mm, 10  10  0.48 mm, and 10  10  0.47 mm, respectively. The samples were weighed before and after they were immersed in the corrosion cell at 23 ± 1 °C. The immersion times tested were 2, 24, or 168 h. Each sample was suspended from nylon strings and completely immersed in 1 l of the solution. The test solution was not renewed during the exposure period and the tests were performed without agitating the solution in the cells. After the experiment, the samples were cleaned by dipping them in a solution of 25% CrO3 for 2 h and then washing them in acetone. The immersion tests were repeated three times to obtain reproducible results. After the immersion tests, X-ray diffraction was performed on the samples to identify the corrosion products that had formed on the surfaces. 3. Results 3.1. Characterization of microstructure Fig. 1a and b shows SEM micrographs of the as-received AZ61 alloy taken on the TD–RD sections. The average size of the grains (d) in the as-received alloy was 15 lm. Irregular-shaped divorced eutectic b-Mg17Al12 phases (Fig. 1b) were observed at grain boundaries. Their area (volume) fraction (as measured using SEM) was 0.01 [9], which is significantly lower than the equilibrium value of 0.06, implying that Al solutes were supersaturated in the matrix. Fig. 2a–c shows the optical micrographs of the as-received,

231

HRDSR 2SP, and HRDSR 3SP samples, respectively, taken from the RD–ND sections. It is evident that extensive shear banding had occurred during HRDSR, resulting in considerable grain refinement and a huge increase in the volume fraction of the b phase (Fig. 2b and c). Numerous very fine b-phase particles were dispersed along the flow lines lying almost parallel to the RD. A SEM micrograph of HRDSR 3SP (Fig. 2d) shows that most of the very fine b-phase particles are located at the boundaries of equiaxed grains. The as-received AZ61 contains a highly supersaturated state of the Al solute atoms, which allows b particles to dynamically precipitate out from the Al supersaturated matrix during deformation by HRDSR. As new grain boundaries that form and evolve from the dislocations generated during shear banding through continuous dynamic recrystallization provide preferential sites for the heterogeneous nucleation of b precipitates as well as fast diffusion paths for the precipitation elements, nucleation, growth and spheroidization of b precipitates preferentially occur at grain boundaries rather than in grains. The area (volume) fractions of the b-phase particles in HRDSR 2SP and HRDSR 3SP, which were measured using SEM, were 0.065 and 0.063, respectively. These values are pretty similar to the equilibrium volume fraction of b phase (0.06) in AZ61, implying that the b phase had fully precipitated during the deformation by HRDSR. Fig. 3a and b shows the TEM micrographs of the as-received AZ61 alloy and HRDSR 3SP, respectively. While the as-received AZ61 alloy has the heterogeneous microstructure composed of mixture of large grains with deformation bands and twins (Fig. 3a) and large grains with low dislocation density, HRDSR 3SP has the fairly homogeneous microstructure composed of very fine grains often with a high density of tangled dislocations. Fig. 3c, d and e, f shows the dark- and bright-field TEM micrographs of HRDSR 2SP and HRDSR 3SP, respectively, taken for better viewing the matrix grain boundaries and b-phase particles in the matrix. Both materials are constituted of fine microstructures composed of very fine equiaxed grains whose average sizes are 1.1 and 1.5 lm, respectively. They have very fine b particles with average sizes of 70 and 140 nm, which are concentrated at grain boundaries. The comparison of the microstructures of the two materials indicates that a higher speed ratio of 3 is more effective in decreasing the sizes of matrix grains and b-phase particles. This may be because the amount of shear strain introduced is larger and the rate of deformation is higher during deformation at the higher speed ratio. Fig. 4 shows the element maps for Al and Mn overlapped on a TEM image of HRDSR 3SP. The green and yellow patches indicate X-ray signals from Al and Mn, respectively. The results of the EDS analysis inform that the majority of the particles are Mg17(Al, Zn)12 and that the yellow minor-phase particles, which are typically smaller, are Al–Mn compounds. Fig. 5a–d shows the optical photographs taken on the TD–RD sections of the unannealed HRDSR 2SP sample and the HRDSR 2SP samples that were annealed at various temperatures for 2 h. The size of the grains increased with increasing annealing temperature. The grains markedly grew at 400 °C. Fig. 6a–e shows the SEM micrographs of the unannealed HRDSR 2SP and the annealed HRDSR 2SP samples. The size of the grains and the volume fractions of the b phase measured using SEM for the annealed samples are listed in Table 1. As with the microstructure of the unannealed sample (Fig. 2d), those of the annealed samples show that most b-phase particles are located at grain boundaries. The volume fractions of the b phase are similar as 0.06–0.07 at all annealing temperatures. The number of b particles whose sizes are in the range 200–500 nm, however, has significantly increased in the microstructures after annealing at temperatures of 240–320 °C because of particle coarsening. In the sample annealed at 400 °C (Fig. 6e), b is found as a continuous phase along grain boundaries, indicating that the b particles in HRDSR 2SP were dissolved into

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(a)

(b)

(c)

(d)

(e)

Fig. 6. SEM micrographs of (a) unannealed HRDSR 2SP sample and HRDSR 2SP samples annealed at (b) 240, (c) 280, (d) 320, or (e) 400 °C for 2 h.

Table 1 Microstructural and electrochemical features of unannealed and annealed samples. The errors in the measurements of Icorr and bc represent the uncertainty in the accuracy of Tafel slopes determined from polarization curves. Material As-received HRDSR 2SP HRDSR 3SP HRDSR 2SP annealed HRDSR 2SP annealed HRDSR 2SP annealed HRDSR 2SP annealed

Ecorr (V vs. SCE)

at at at at

240 °C 280 °C 320 °C 400 °C

1.49 1.49 1.47 1.48 1.5 1.49 1.48

Icorr (lA/cm2)

bc (mV/dec

17.71 ± 3.66 10.48 ± 3.97 6.45 ± 2.12 14.14 ± 2.14 13.04 ± 2.88 5.96 ± 1.98 16.06 ± 3.79

142.96 ± 10.53 133.82 ± 13.21 135.79 ± 17.18 145.38 ± 8.53 135.10 ± 6.77 129.81 ± 10.51 147.86 ± 13.49

matrix and then reprecipitated at grain boundaries during the annealing process. Fig. 7a–c shows the (0 0 0 2) pole figures of the as-received, HRDSR 2SP, and HRDSR 3SP samples taken at s = 9. The maximum intensity of the basal texture sharply increased from 21 to 50–59 during HRDSR. Basal texture weakening, which reportedly occurs in Mg alloys during differential speed rolling [20,21], did not occur in this study. This may be because of the processing temperature used, which was chosen to be relatively low in order to maximize the efficiency of grain refinement. The basal texture typically strengthens in Mg alloys at low deformation temperatures because basal slip and tension twinning dominantly occur at low temperatures. Fig. 7d–f shows the (0 0 0 2) pole figures of the HRDSR 2SP samples annealed at various temperatures. The basal texture was retained during annealing. The maximum intensity of the basal texture did not significantly change in the samples annealed up to 280 °C. The intensity of the texture notably decreased, however, in the samples annealed at 320 and 400 °C. It appears that the grains oriented far from the basal texture component grew rapidly

1

)

Grain size (lm)

Vol. fraction of b phase

17.4 1.5 1.09 2.7 4.3 5.4 13.1

0.01 0.065 0.063 0.056 0.078 0.077 0.072

during annealing above 320 °C, resulting in the basal texture weakening. 3.2. Corrosion measurements 3.2.1. Polarization measurements Fig. 8a shows the potentiostatic polarization curves of the as-received, HRDSR 2P, and HRDSR 3SP samples. All three materials show self-passivation characteristics because passivity plateaus exist on the anodic sides of the polarization curves. The breakdown potentials of the HRDSR-processed samples are nobler than that of the as-received sample, implying that the passivity of the film surface increases with decreasing grain size. The values of corrosion potential (Ecorr) and corrosion current density (Icorr) derived from the cathodic curves are listed in Table 1. Ecorr for all three materials ranged from 1.47 to 1.49 VSCE. The Icorr for the HRDSR-processed AZ61 alloys is lower than that for the as-received sample, suggesting that the grain-size reduction by HRDSR decreases the rate of corrosion of the alloys. According to the polarization curves, the

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As-received

(a)

HRDSR 2SP

HRDSR 3SP

TD

RD

(b)

Fig. 7. (0 0 0 2) pole figures of (a) as-received, HRDSR 2SP, and HRDSR 3SP samples. (0 0 0 2) pole figures of HRDSR 2SP samples annealed at 280, 320, or 400 °C for 2 h.

(b) -0.6 -0.7 -0.8 -0.9 -1.0 -1.1 -1.2 -1.3 -1.4 -1.5 -1.6 -1.7 -1.8 -1.9 -2.0 -8 10

As-received HRDSR 2SP HRDSR 3SP

Potential, V(vs. SCE)

Potential, V(vs. SCE)

(a)

10

-7

10

-6

10

-5

10

-4

10

-3

10

-2

Current density (A/cm2)

-0.6 -0.7 -0.8 -0.9 -1.0 -1.1 -1.2 -1.3 -1.4 -1.5 -1.6 -1.7 -1.8 -1.9 -2.0 -7 10

HRDSR 2SP o annealed at 240 C o annealed at 280 C o annealed at 320 C o annealed at 400 C

10

-6

10

-5

10

-4

10

-3

10

-2

Current density (A/cm2)

Fig. 8. (a) Potentiostatic polarization curves for as-received, HRDSR 2P, and HRDSR 3SP samples. (b) Potentiostatic polarization curves for HRDSR 2P samples annealed at various temperatures.

decrease of corrosion rate results from simultaneous decrease in anodic and cathodic kinetics. Fig. 8b shows the potentiostatic polarization curves for the HRDSR 2SP samples annealed at various temperatures. There is no significant difference in anodic and cathodic reaction kinetics for the samples annealed up to 280 °C. Icorr considerably decreases for the sample annealed at 320 °C (Table 1) as a result of simultaneous decrease in the rates of the anodic and cathodic reactions. However, Icorr increases again for the sample annealed at 400 °C. 3.2.2. Weight-loss measurements The rate of corrosion calculated from the weight loss of the sample due to corrosion provides the most straightforward measurement of corrosion resistance. Fig. 9a shows the results of the weight loss test for the as-received, HRDSR 2SP, and HRDSR 3SP samples immersed in the corrosion cell for 168 h. The rate of weight loss tends to decrease as grain size decreases. The HRDSR 3SP sample exhibits a 38.9% lower rate of weight loss than the as-received sample. Annealing the HRDSR-processed samples sig-

nificantly affects the rate of weight loss, as shown in Fig. 9b. Unlike the trend shown in Fig. 9a where the rate of weight loss increases with increasing grain size, the rate of weight loss decreases with increasing grain size in the samples annealed up to 320 °C and then increases at 400 °C with further increasing grain size. The HRDSR 2SP sample annealed at 320 °C exhibits the lowest corrosion rate, a 24% lower rate of weight loss than does the unannealed HRDSR 2SP sample. These results observed in Fig. 9a and b are in agreement with those observed in the measurements of Icorr from the polarization curves (Fig. 8a and b). Fig. 9c shows the rates of weight loss of the as-received, HRDSR 2SP, and HRDSR 2SP samples annealed at 320 °C for 2 h as a function of immersion time. The results show that grain refinement by HRDSR and post-HRDSR annealing decrease the rate of corrosion from the early stage of immersion. The rate of corrosion is, however, relatively fast compared to that for the long-term immersions. XRD analysis was used to determine the corrosion products that had formed on the surfaces of the HRDSR 2SP sample immersed for 168 h (Fig. 10). The exposed surfaces were directly analysed

H.S. Kim, W.J. Kim / Corrosion Science 75 (2013) 228–238 18

18

Immersion time : 168 h

Weight loss rate Grain size

16

(b)

16 14

12

12

10

10

8

8

6

6

4

4

2

2 0

0

As-received

HRDSR 2SP

14

16

HRDSR 2SP

14

10

10

8

8 6 6 4

4 Weight loss rate Grain size

2 0 No Annealing 240

HRDSR 3SP

2 0

280

320

400

Annealing temperature (oC)

Material type 100

Weight Loss Rate (μg/cm2h)

(c)

12

Immersion time : 168 h 12

Grain size (μm)

14

Grain size (μm)

Weight Loss Rate (μg/cm2h)

(a)

Weight Loss Rate (μg/cm2h)

234

As-receieved HRDSR 2SP o HRDSR 2SP annealed at 320 C

10

3 1

10

100

400

Immersion time (h) Fig. 9. (a) Results of weight loss tests for as-received, HRDSR 2SP, and HRDSR 3SP samples immersed for 168 h. (b) Results of weight loss tests for the HRDSR 2SP samples annealed at various temperatures (immersed for 168 h). (c) Results of weight loss tests for as-received, HRDSR 3SP, and HRDSR 2SP samples annealed at 320 °C for 2 h immersed for different time durations. Error bars represent ± standard error of the mean values of the experimental results.

Mg(OH)2 film due to its dehydration during the sample preparation for SEM observation. The size and number of the Mg(OH)2 islands decreased as the annealing temperature increased, indicating that the surface corroded less and that the corrosion was more uniform as the higher annealing temperature was used.

Intensity / a. u.

Mg Mg17 Al12 Mg(OH)2

4. Discussion 4.1. Grain size

20

30

40

50

60

70

80

2θ Fig. 10. XRD analysis of phases on surface of HRDSR 2SP immersed for 168 h.

without cleaning the corrosion products off the surfaces. The analysis results indicates that the corrosion products mainly consisted of Mg(OH)2. Fig. 11a–d shows the surfaces of the as-received sample, the unannealed HRDSR 2SP sample, and the annealed HRDSR 2SP samples (at 280 or 320 °C) after immersion of 2 h. The corroded surface of the as-received alloy exhibits typical microgalvanic corrosion features. The a phase, which is adjacent to the divorced eutectic (i.e. coarse b) phase located at grain boundaries, preferentially corroded. This type of corrosion was rarely observed on the surfaces of the unannealed HRDSR 2SP and annealed HRDSR 2SP samples. Instead, Mg(OH)2 islands associated with cracks were often encountered. Those cracks most likely originated from shrinkage of thick

Ralston et al. [22] showed that the rate of passive corrosion of metals is proportional to the reciprocal square root of grain size. Icorr and the rates of weight loss obtained from the potentiodynamic polarization and immersion tests, respectively, for the as-received, HRDSR 2SP, and HRDSR 3SP samples are plotted against the reciprocal square root of the average grain size (d 1/2) in Fig. 12a and b, respectively. The linear relationship is obtained from Icorr and rate of weight loss that decrease with increasing d 1/2. However, the linear relations do not hold when the data for the annealed HRDSR 2SP samples are considered together. The annealed samples, especially those annealed at 280 and 320 °C, exhibit lower rates of corrosion than expected from the linear relation (Fig. 12b). This result implies that the corrosion resistance of AZ61 alloy also depends on structural factors other than grain size. There have been several reports that the rate of corrosion decreases when the size of grains in Mg alloys decreases [23–25]. Song et al. [26,27], however, reported contradictory results. It has been suggested that grain boundaries can act as a physical barrier against corrosion by facilitating passivation if the corrosion product formed on surface of the sample is stable [28]; otherwise, grain boundaries accelerate the rate of corrosion as they function as local anodes. The corrosion product Mg(OH)2 forms on

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(b)

(a) β

(c)

(d)

Fig. 11. Surfaces of (a) as-received sample, (b) HRDSR 2SP sample, and HRDSR 2SP samples annealed at either (c) 280 or (d) 320 °C after immersion for 2 h.

(a) 24

(b) 20 Weight Loss Rate (μg/cm2h)

22 20

Icorr (μA/cm2)

18 16 o

14

400 C

12

o

240 C o

280 C

10 8 6

o

320 C

4 2 0 0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

1.2

d-1/2 (μm-1/2)

18 16 14

o

400 C 12 10

o

280 C 8

o

320 C

6 4 2 0 0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

1.2

d-1/2 (μm-1/2)

Fig. 12. (a) Icorr and (b) rates of weight loss obtained from potentiodynamic polarization and immersion tests, respectively, of as-received, HRDSR 2SP, HRDSR 3SP, and annealed HRDSR 2SP samples, which are plotted against d 1/2. Error bars represent ± standard error of the mean values of the experimental results.

a MgO layer on the surface of Mg substrates during the immersion. Thus, the stability of the MgO layer on the surface of the Mg substrate is important; if the MgO layer is less prone to cracking, breaking of the Mg(OH)2 layer may be subsequently prevented. There is free volume mismatch between MgO layer and Mg substrate because of the difference in their molar volumes, so when MgO layer form on Mg substrate, tensile stress builds up in MgO layer [29], increasing the susceptibility of cracking of MgO layer or interface between MgO and Mg. Although it has been proposed that a fine-grained microstructure provides a means of relieving the stress through grain-boundary diffusion of vacancies, thus decreasing the degree of cracking in the MgO layer [29], there has not been any evidence yet to support this proposal. We therefore offer an alternative proposal for the enhanced stability of the MgO layer on the Mg substrate composed of small grains. Fig. 13a shows a TEM image of HRDSR 3SP, showing that a dense nanocrys-

talline MgO layer is deposited on the surface of HRDSR 3SP. High resolution TEM image and electron diffraction pattern were taken in the region just below the MgO layer (Fig. 13b). The inverse fast Fourier transform (IFFT) of the image in (b) using the selected spots in the diffraction pattern ring in (b) reveals that isolated MgO are surrounded by Mg matrix (Fig. 13c). These MgO nanocrystals may have been formed as oxygen atoms diffused from the surface into the grain boundaries of the matrix. If this is so, the distribution of MgO in the Mg matrix below the MgO layer will be more uniform and the volume fraction of MgO nanocrystals will be higher when grain boundary area per unit volume is larger (i.e. when grain size is smaller). A schematic for this situation is shown in Fig. 14. The formation of a layer containing a mixture of MgO and Mg phases between the MgO layer and the Mg substrate should be beneficial in decreasing the amount of tensile stress in the MgO layer by acting as a buffer layer that decreases the

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(a)

MgO

100 nm

(b)

(c)

MgO

Mg

Fig. 13. (a) A TEM image taken around the surface region of HRDSR 3SP. (b) High resolution TEM image of HRDSR 3SP taken in the Mg substrate under MgO layer (a boxed area in (a)). (c) The inverse fast Fourier transform (IFFT) of the image in (b) using the selected spots in the diffraction pattern ring in (b).

MgO MgO + Mg Mg

Fine-grained Mg MgO

Mg

Coarse-grained Mg Fig. 14. Schematic for formation of intermediate layer composed of mixture of Mg and MgO phases between Mg substrate and MgO layer.

to always increase the rate of anodic dissolution rate [30,31]. For this reason, the corrosion resistance of metals is enhanced when the density of dislocations in the matrix is decreased through recovery. Song et al. [26] reported that annealing ECAP-processed Mg below recrystallization temperature was beneficial for increasing its resistance to corrosion. In this study, the corrosion rate decreased as the annealing temperature increased up to 320 °C (Fig. 9b), suggesting that density of dislocations decreased with increasing temperature. Generally, annealing is accompanied by grain growth. Under proper annealing condition, however, the positive effect of the decrease in the density of dislocations through recovery can be dominant over the negative effect of the increase in the size of grains by grain growth. This condition appears to have been optimally met for the sample annealed at 320 °C. When significant grain growth takes place as in the sample annealed at 400 °C, however, the corrosion resistance may decrease due to the pronounced negative effect of the increase in grain size. 4.3. Texture

sharpness of the stress gradient across the boundary between the MgO layer and the Mg substrate. 4.2. Dislocation density Although the effect of grain boundaries on the rate of corrosion is controversial, a high density of dislocations in the matrix seems

It has previously been shown that the rate of corrosion of the AZ31 alloy decreases with increasing intensity of the (0 0 0 1) texture [13,14]. This is because the (0 0 0 1) surface has the lowest surface energy and thus dissolves slower than the (1 0 1 0) and (1 1 2 0) surfaces in Mg. In this study, the intensity of the basal texture significantly increased after HRDSR (Fig. 7). Thus, this basal texture

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strengthening is possible to have contributed to the enhanced corrosion resistance of the AZ61 alloy. At the annealing temperature of 320 °C, however, the best corrosion resistance was obtained despite the sharp decrease of the maximum intensity of basal texture after annealing. This result indicates that the texture effect on the corrosion resistance is not so strong compared to the effect of dislocation density. 4.4. Presence of the b phase The b-phase is known to play two contradicting roles in the corrosion of Mg–Al alloys. It acts as a cathode in galvanic action, leading to accelerated corrosion of a-Mg matrix near b-phase [32]. However, it can also act as a barrier against the propagation of corrosion in the a-Mg matrix, because the b phase is much inerter than the surrounding matrix in chemical activity [33], if fraction of b phase high and it is continuous like a net over the a-Mg matrix by being distributed between the dendritic arms or along grain boundaries [34]. Song et al. [27] compared the corrosion behaviours of ECAP-processed AZ91 alloy with small grains (2 lm) and as-cast AZ91 alloy with coarse grains (50 lm). They found that the ECAP-processed alloy exhibited decreased resistance to corrosion and attributed the result partially to the morphology change of b phase from the net-like structure into a large number of particles during the ECAP process. Unlike HRDSR 2SP and HRDSR 3SP, however, the ECAP-processed AZ91 alloy studied by Song et al. [27] contained a considerable amount of relatively coarse b-phase particles (10 lm). Argade et al. [16] showed that the susceptibility of microgalvanic corrosion could be considerably lowered upon refinement of a secondary phase to nanosized particles. They reported that the friction-stir-processed Mg–Y–Re alloy containing nanoparticles (<100 nm) exhibited significantly lower galvanic corrosion activity compared to the same composition alloy containing coarse particles. This finding is in agreement with the current observation that the HRDSR-processed AZ61 alloys with nanosized particles exhibit the low activity of microgalvanic corrosion. A high volume fraction of nanosized b-phase particles that are uniformly distributed on grain boundaries of ultrafine grains may also act as a barrier against corrosion. This is possible because the interspacing between the b-phase particles is very small. In the sample annealed at 400 °C, b is present as continuous phase along grain boundaries, but its role as a barrier against the propagation of corrosion is considered to be limited because the distance between the b phase is large such that a phase is not effectively protected by the b phase. Rather, accelerated corrosion due to microgalvanic corrosion is expected to occur. 5. Conclusions The corrosion behaviour of the ultrafine-grained AZ61 alloys containing very fine b-phase particles, which were prepared using the HRDSR technique, was investigated using immersion tests and electrochemical analysis. The effect of annealing on the resistance to corrosion was also examined. The results can be summarized as follows: (1) The results of both the immersion and electrochemical tests showed that decreasing the size of grains was beneficial in improving the corrosion resistance of the alloy. The improved corrosion resistance with finer grains was attributed to the enhanced passivity of the oxide film that formed on the surface of the alloy. The formation of a many isolated MgO nanocrystals in the fine-grained Mg substrate below the MgO layer, which may have been formed as oxygen atoms diffused from the surface into the grain boundaries of the matrix, was suggested to enhance the stability of

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the MgO layer by providing a buffer layer that decreases the sharpness of the stress gradient across the boundary between the MgO layer and the Mg substrate. (2) The rate of corrosion measured from the electrochemical and immersion tests of the AZ61 alloy has a linear relation with the reciprocal square root of the average grain size. Annealing increased the corrosion resistance of the alloy more than expected from the linear relation, because annealing decreased the density of dislocations in the matrix. When extensive grain growth occurred during annealing, however, the corrosion resistance was decreased because of the pronounced negative effect of the increase in grain size. (3) Very fine b particles preferentially located at grain boundaries of the HRDSR-processed AZ61 alloys did not show evidence of galvanic corrosion; however, coarse b particles showed such evidence, indicating that refining the b phase to nanosized particles decreases the susceptibility of individual microgalvanic corrosion between the a-Mg matrix and the b phase. (4) HRDSR offers an attractive possibility for simultaneously improving the corrosion properties and mechanical properties (including room-temperature strength and superplasticity) of AZ61 alloy sheets.

Acknowledgement This research was supported by Basic Science Research Program through the National Research Foundation of Korea (2013) funded by the Ministry of Education, Science and Technology (NRF2013R1A1A2010637). References [1] H. Jones, E.J. Lavernia, in: M.A. Otooni (Ed.), Elements of Rapid Solidification, Springer, 1998, p. 135. [2] W.J. Kim, S.I. Hong, Y.S. Kim, S.H. Min, H.T. Jeong, J.D. Lee, Texture development and its effect on mechanical properties of an AZ61 Mg alloy fabricated by equal channel angular pressing, Acta Mater. 51 (2003) 3293–3307. [3] A. Yamashita, Z. Horita, T.G. Langdon, Improving the mechanical properties of magnesium and a magnesium alloy through severe plastic deformation, Mater. Sci. Eng. A300 (2001) 142–147. [4] S.R. Agnew, J.A. Horton, T.M. Lillo, D.W. Brown, Enhanced ductility in strongly textured magnesium produced by equal channel angular processing, Scripta Mater. 50 (2004) 377–381. [5] Y. Yoshida, L. Cisar, S. Kamado, Y. Kojima, Effect of microstructural factors on tensile properties of an ECAE-processed AZ31 magnesium alloy: special issue on platform science and technology for advanced magnesium alloy, II, Mater. Trans. 44 (2003) 468–475. [6] K. Matsubara, Y. Miyahara, Z. Horita, T.G. Langdon, Developing superplasticity in a magnesium alloy through a combination of extrusion and ECAP, Acta Mater. 51 (2003) 3073–3084. [7] M. Mabuchi, H. Iwasaki, K. Yanase, K. Higashi, Low temperature superplasticity in an AZ91 magnesium alloy processed by ECAE, Scripta Mater. 36 (1997) 681– 686. [8] W.J. Kim, I.B. Park, Enhanced superplasticity and diffusional creep in ultrafinegrained Mg–6Al–1Zn alloy with high thermal stability, Scripta Mater. 68 (2013) 179–182. [9] W.J. Kim, I.B. Park, S.H. Han, Formation of a nanocomposite-like microstructure in Mg–6Al–1Zn alloy, Scripta Mater. 66 (2012) 590–593. [10] W.J. Kim, B.G. Hwang, M.J. Lee, Y.B. Park, Effect of speed-ratio on microstructure, and mechanical properties of Mg–3Al–1Zn alloy, J. Alloys Compd. 509 (2011) 8510–8517. [11] G. Song, A. Atrens, X. Wu, B. Zhang, Corrosion behaviour of AZ21, AZ501 and AZ91 in sodium chloride, Corros. Sci. 40 (1998) 1769–1791. [12] R. Ambat, N.N. Aung, W. Zhou, Evaluation of microstructural effects on corrosion behaviour of AZ91D magnesium alloy, Corros. Sci. 42 (2000) 1433– 1455. [13] R. Xin, Y. Luo, A. Zuo, J. Gao, Q. Liu, Texture effect on corrosion behaviour of AZ31 Mg alloy in simulated physiological environment, Mater. Lett. 72 (2012) 1–4. [14] G.L. Song, R. Mishra, Z.Q. Xu, Crystallographic orientation and electrochemical activity of AZ31 Mg alloy, Electrochem. Commun. 12 (2010) 1009–1012.

238

H.S. Kim, W.J. Kim / Corrosion Science 75 (2013) 228–238

[15] Ji. Liao, M. Hotta, N. Yamamoto, Corrosion behaviour of fine-grained AZ31B magnesium alloy, Corros. Sci. 61 (2012) 208–214. [16] G.R. Argade, S.K. Panigrahi, R.S. Mishra, Effects of grain size on the corrosion resistance of wrought magnesium alloys containing neodymium, Corros. Sci. 58 (2012) 145–151. [17] Z. Pu, S. Yang, G.L. Song, O.W. Dillon Jr., D.A. Puleo, I.S. Jawahir, Ultrafinegrained surface layer on Mg–Al–Zn alloy produced by cryogenic burnishing for enhanced corrosion resistance, Scripta Mater. 65 (2011) 520–523. [18] D. Song, A. Ma, J. Jiang, P. Lin, D. Yang, J. Fan, Corrosion behaviour of equalchannel-angular-pressed pure magnesium in NaCl aqueous solution, Corros. Sci. 52 (2010) 481–490. [19] M. Alvarez-Lopez, M.D. Pereda, J.A. del Valle, M. Fernandez-Lorenzo, M.C. Garcia-Alonso, O.A. Ruano, M.L. Escudero, Corrosion behaviour of AZ31 magnesium alloy with different grain sizes in simulated biological fluids, Acta Biomater. 6 (2010) 1763–1771. [20] X. Huang, K. Suzuki, A. Watazu, I. Shigemastu, N. Saito, Improvement of formability of Mg–Al–Zn alloy sheet at low temperatures using differential speed rolling, J. Alloys Compd. 470 (2009) 263–268. [21] W.J. Kim, J.B. Lee, W.Y. Kim, H.T. Jeong, H.G. Jeong, Microstructure and mechanical properties of Mg–Al–Zn alloy sheets severely deformed by asymmetrical rolling, Scripta Mater. 56 (2007) 309–312. [22] K.D. Ralston, N. Birbilis, C.H.J. Davies, Revealing the relationship between grain size and corrosion rate of metals, Scripta Mater. 63 (2010) 1201–1204. [23] N.N. Aung, W. Zhou, Effect of grain size and twins on corrosion behaviour of AZ31B magnesium alloy, Corros. Sci. 52 (2010) 589–594. [24] T. Zhang, Z. Ji, S. Wu, Effect of extrusion ratio on mechanical and corrosion properties of AZ31B alloys prepared by a solid recycling process, Mater. Des. 32 (2011) 2742–2748.

[25] L. Jinsun, H. Makoto, Y. Naotsugu, Corrosion behavior of fine-grained AZ31B magnesium alloy, Corros. Sci. 61 (2012) 208–214. [26] D. Song, A.B. Ma, J.H. Jiang, P.H. Lin, D.H. Yang, J.F. Fan, Corrosion behaviour of equal-channel-angular-pressed pure magnesium in NaCl aqueous solution, Corros. Sci. 52 (2010) 481–490. [27] D. Song, A.B. Ma, J.H. Jiang, P.H. Lin, D.H. Yang, J.F. Fan, Corrosion behaviour of bulk ultra-fine grained AZ91D magnesium alloy fabricated by equal-channel angular pressing, Corros. Sci. 53 (2011) 362–373. [28] L. Liu, Y. Li, F. Wang, Pitting mechanism on an austenite stainless steel nanocrystalline coating investigated by electrochemical noise and in-situ AFM analysis, Electrochim. Acta 54 (2008) 768–780. [29] D. Orlov, K.D. Ralston, N. Birbilis, Y. Estrin, Enhanced corrosion resistance of Mg alloy ZK60 after processing by integrated extrusion and equal channel angular pressing, Acta Mater. 59 (2011) 6176–6186. [30] M. Ben-Haroush, G. Ben-Hamu, D. Eliezer, L. Wagner, The relation between microstructure and corrosion behaviour of AZ80 Mg alloy following different extrusion temperatures, Corros. Sci. 50 (2008) 1766–1778. [31] G. Ben-Hamu, A. Eliezer, E.M. Gutman, Electrochemical behaviour of magnesium alloys strained in buffer solutions, Electrochim. Acta 52 (2006) 304–313. [32] K. Nisancioglu, O. Lunder, T.K. Aune, Corrosion mechanism of AZ91 magnesium alloy, in: Proceedings of 47th World Magnesium Association, Mcleen, Virginia, 1990, p. 43. [33] W.S. Loose, in: L.M. Pidgeon, J.C. Mathes, N.E. Woldmen (Eds.), Corrosion and Protection of Magnesium, ASM International, Materials Park, OH, 1946, p. 173. [34] G. Song, A. Atrens, M. Dargusch, Influence of microstructure on the corrosion of diecast AZ91D, Corros. Sci. 41 (1999) 249–273.