Effect of strain rate on acoustic emission during hydrogen assisted cracking in high carbon steel

Materials Science and Engineering A 550 (2012) 408–417

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Effect of strain rate on acoustic emission during hydrogen assisted cracking in high carbon steel E.D. Merson, M.M. Krishtal, D.L. Merson, A.A. Eremichev, A. Vinogradov ∗,1 Laboratory for the Physics of Strength of Materials and Intelligent Diagnostic Systems, Togliatti State University, Togliatti 445667, Russia

a r t i c l e

i n f o

Article history: Received 2 December 2011 Received in revised form 26 April 2012 Accepted 26 April 2012 Available online 3 May 2012 Keywords: Hydrogen assisted cracking Acoustic emission Strain rate High carbon steel

a b s t r a c t Hydrogen charging of a high carbon spring steel exerts a profound effect on the ductility and damage accumulation during three-point-bending testing. Since the acoustic emission (AE) technique reflects perhaps in the best way microfracturing dynamics, it is found particularly useful for real-time capturing of hydrogen assisted cracking and overall in situ tracing of damage evolution during deformation. In the present work the propensity for hydrogen assisted cracking, its mechanisms and strain rate dependence are investigated using integral characterization of AE paired with microstructure investigations. Analysis of experimental data suggests that two competitive mechanisms operate during three-point bending deformation of the high carbon steel: (i) highly localized plastic deformation and (ii) brittle intercrystalline and transcrystalline cracking promoted by hydrogen influence. The synergistic interplay between these two mechanisms is discussed. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Hydrogen assisted cracking (HAC, referred sometimes historically to as “hydrogen embrittlement”) is a well-known phenomenon causing substantial degradation of properties in steels and high strength alloys. Particularly, the deteriorative influence of hydrogen is essential for ductility, fracture toughness, fatigue crack growth resistance, etc. [1]. Even a very small amount of dissolved hydrogen, on the level lessthan one part per million, can cause a severe detrimental effect on steel performance [2]. Being the root cause of many service-related failures and technological catastrophes HAC attracted much attention from material scientists, chemists, physicists and engineers aiming at failure risk reduction and alleviation of the damage effects induced by hydrogen. In fact, HAC is a generic phenomenon that is relevant to many common failure mechanisms involved in electroplate embrittlement [3], slow strain rate embrittlement [4,5], stress corrosion cracking [6–8], delayed fracture [2,9], weld heat affected zone cracking [10–12], etc. A comprehensive review of HAC and relevant phenomena is given in [13–15]. There is a general agreement in literature that the hydrogen induced effects in steels cannot be described by a single universal mechanism [13]. Robertson and Birnbaum [16] have concluded that among many mechanisms proposed to explain the observed effects of hydrogen on deformation

∗ Corresponding author. Tel.: +7 8482 546303. E-mail addresses: [email protected], [email protected] (A. Vinogradov). 1 On leave from Osaka City University, Osaka 558-8585, Japan. 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.04.094

and fracture, only three mechanisms are preferable: the stressinduced formation of hydrides and their subsequent cleavage, the hydrogen-enhanced localized plasticity – the HELP mechanism, and the decohesion mechanism – hydrogen reduction of the cohesive strength of the lattice. However, the significant hydrogen embrittlement is often observed at so low average amount of hydrogen that none of the above mechanisms is capable of explaining it. To enforce these mechanisms it is essential that a sufficient amount of hydrogen atoms should be aggregated locally. This can be achieved either by diffusion or dislocation transport. Hence, the features of these mechanisms govern the specific temperature and strain rate dependence of HAC in materials. Bastien and Azou first pointed out the negative strain rate dependence of hydrogen embrittlement in steel in 1951 [17]. They found that the ductility of a hydrogen-charged steel decreases with the decrease in the strain rate at constant temperature. Later Brown and Baldwin [18] investigated the details of this dependence. They have shown that HAC is particularly pronounced at relatively low strain rates at room temperature and is virtually absent at high strain rates. The significance of this finding is twofold: (i) estimating the materials propensity to HAC requires tedious tests which are time-consuming and costly because the slow strain rates are required to reveal the features of HAC; (ii) the most favorable conditions for development of HAC – ambient temperature and low strain rates – correspond to the operating conditions of the vast majority of engineering structures and components. Based on the foregoing, it can be concluded that the management of industrial facilities, timely maintenance and eventually prevention of hydrogen-affected technological catastrophes relies heavily on in depth understanding of the HAC phenomenon on one

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Table 1 Chemical composition of steel 70. C

Si

Mn

Cr

S

P

Cu

Ni

Fe

0.70

0.29

0.72

0.03

0.006

0.013

0.06

0.02

Balance

Table 2 Hydrogen content, maximum deflection max and cumulative AE count N in the high carbon steel 70. Specimen state

CH , cm3 /100 g

˙ s−1 Initial strain rate ε, 3 × 10−3

Reference Hydrogen charged Dehydrogenated

1.1 12.4 9.0

3 × 10−4

3 × 10−6

max , mm

N

max , mm

N

max , mm

N

18.3 14.6 18.4

32 84 32

17.3 10.1 16.8

60 385 36

16.9 6.5 11.0

317 5172 437

hand and on the success in development of time-efficient testing schemes backed by a reliable technique capable of real time monitoring of HAC in loaded metallic articles on the other. In response to this demand a massive literature on the hydrogen induced effects in metals and non-destructive testing (NDT) methods has evolved [19–21]. In spite of all efforts, hydrogen assisted fracture remains to be a primary cause of premature failure of high strength metals, whereas no commonly accepted NDT technology for early identification of the hydrogen-induced damage has emerged to meet industrial standards. Acoustic emission (AE) is among the most promising techniques for this purpose. The method is featured by its unique capacity of delivering real time information about processes of structural rearrangements within a solid body which is a necessary prerequisite for the monitoring of damage accumulation and fracture kinetics. Acoustic emissions are the transient elastic waves generated by the rapid release of energy within a material under external influence. The dislocation-mediated plastic deformation and the crack nucleation and advance have long been recognized as primary AE sources in materials under load [22]. The former generates usually a sort of continuous noise-like signal due to largely overlapping small transients created by intermittent dislocation motion. The latter generates AE bursts of relatively high amplitudes.The high sensitivity of AE technique to HAC has been convincingly demonstrated in a large number of publications, e.g. [19,21,23–25]. A vast majority of AE tests associated with HAC have been performed so far using either constant load (stress) conditions in stress corrosion cracking (SSC) testing [26,27], slow strain rate tests (SSR, with nominal extension rate in the range of 1 × 10−7 –1 × 10−5 s−1 ) utilizing smooth samples or fracture toughness and crack propagation tests on notched and/or pre-cracked samples [28]. No AE studies concerning the strain rate dependence of hydrogen induced damage accumulation have been performed so far to the author’s best knowledge. Understanding of the interaction between the mechanisms involved into HAC under different strain rates is supposed to shed some light on the mutual role of these mechanisms and their kinetics on different stages of fracture. Thus, the objective of this work is to clarify the effect of hydrogen ingress on interplay between plastic deformation and fracture in cathodically charged high carbon steel with a help of AE technique in dependence on the strain rate.

2. Experimental details High carbon steel 70 in Russian designation (AISI 1070 is a close analogue) was chosen for the present study because it is widespread industrial use as spring steel, which is known to be prone to premature failure due to HAC [29]. The chemical composition of this steel is shown in Table 1. The samples were machined to have a rectangular shape of 120 mm length, 20 mm width and 1.8 mm

thickness. They were thermally treated using a schedule that is commonly used for this type steels: (1) quenching from 850 ± 10 ◦ C in oil; (2) low tempering at temperature of 240 ± 10 ◦ C for 1 h; (3) primary tempering at 450 ± 10 ◦ C for 2 h followed by cooling in air. The hardness after this treatment was of 49 ± 2 HRC. For hydrogen charging a part of the specimen was electrolitically plated with a Zn layer of 10 ␮m thick from an alkaline solution containing of 110–120 g/l NaOH and 12–14 g/l ZnO. The plating was performed at electric current density of 8 A/dm2 during 20 min. Reference specimens were left unplated for comparison. A part of Zn-plated specimens was subjected to conventional dehydrogenation heat treatment (DHT) by annealing at 190 ◦ C for 3 h. The hydrogen content CH was measured in all specimens by a LECO RH402 hydrogen analyzer with a nominal accuracy less than 2 ppm, Table 2. Three points bending mechanical testing was performed at room temperature using a screw driven Instron-type frame. The load was applied at three different traverse velocities resulting in three different initial strain rates: ‘slow’ – ε˙ = 3 × 10−6 s−1 , ‘middle’ 3 × 10−4 s−1 and ‘high’ 3 × 10−3 s−1 . The slowest strain rate corresponds to the upper bond of the strain rates used typically during SSRT. The ductility was characterized by the maximum deflection at break max . Acoustic emission was recorded using a home-made PCcontrolled system with a 12 bits ADC at the core. A broadband (50–500 kHz) piezoelectric transducer MSAE-L2 (Microsensors AE Ltd., Russia) with a low noise built-in 27 dB preamplifier was securely mounted on the specimen using vacuum oil as a couplant as shown in the schematic setup in Fig. 1. The signal from the sensor’s preamp output was transferred through a main filteramplifier MSAE-FA010 with the gain set at 40 dB and the frequency band set between 50 and 1200 kHz. The laboratory noise did not exceed 30 ␮V (peak-to-peak). The burst AE signals having relatively high amplitude above 1 mV (of 30 dB threshold) were recorded and counted synchronously with the load signal. The waveforms of 4096 readings were stored for the post-mortem analysis: the energy per realization was calculated from the power spectral density function as described in detail in [30]. The MSAE-FA010 amplifier contains a built-in circuit allowing precise measurement of the true root-mean-square voltage URMS with 100 ms integration time. The URMS signal was fed to a 14 bit ADC for continuous acquisition at 1 kHz sampling rate. The fracture surface was observed by a scanning electron microscope LEO1455VP.

3. Results Electroplating results in steep increase of hydrogen content (Table 2). Despite the fact that the conventional dehydrogenation procedure by annealing at 190 ◦ C for 3 h reduces the CH value,

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Fig. 1. Schematics of experimental setup.

the latter remains relatively large. After hydrogenation by Znelectroplating the ductility in bending max notably reduces at all strain rates, including the highest one (cf. Table 2). The effect of embrittlement is more pronounced at slower strain rates, as is reasonably expected from a large volume of experimental data reported in the literature. After DHT the ductility recovers at relatively high strain rates; however at the slow strain rate of 3 × 10−6 s−1 it is still low. The scanning micrographs of the fracture surfaces reveal fracture patterns showing mixed ductile–brittle features, Figs. 2 and 3. The fraction of brittle and ductile components of the relief varies strongly depending on testing conditions. Commonly, the decrease in ε˙ and/or the increase of hydrogen concentration CH gives rise to the increasing fraction of brittle facets in the fracture relief. Indeed, the fracture surface of the reference specimens tested at the highest strain rate of 3 × 10−3 s−1 is almost entirely ductile, Fig. 2a. The ductile transcrystalline microcracks can also be seen in a central part of fracture surface as typical “steps” between the MnS sulfides, Fig. 2a.The X-ray energy-dispersive analysis, Fig. 2c, revealed an excess amount of S and Mn atoms in the local area corresponding to the steps clearly visible at the fracture relief. Ductile dimples between such inclusions (showed by arrows in Fig. 2b) serve as a clear evidence for localized plastic deformation accompanying coalescence of voids surrounding the inclusions. The fracture surface appearance of coated specimens is distinctly different. This is particularly clear for the specimens tested at the slow strain rate 3 × 10−6 s−1 . Fig. 3a reveals a brittle zone having an approximately elliptic shape in the central part of fracture surface. This zone extends across the whole section of the specimen (see the red marked area). One can notice a macrocrack aligned with the major ellipse axis in the middle part of the fracture surface. This crack extends in the direction normal to the fracture surface and propagates in the intercrystalline manner across the whole brittle zone marked by B, although the signatures of sulfides decohesion, such as those seen in Fig. 2, are rare, Fig. 3b and c. Fracture surface patterns of other specimens, including DHT ones, are of mixed type between the above discussed two extremes illustrated in Figs. 2 and 3. Hence, one can see that the area fraction of ductile or brittle component is controlled by the strain rate and the amount of dissolved hydrogen. Plastic deformation due to dislocation glide in single phase pure materials produces a continuous noise-like acoustic emission signal [22]. Hence, continuous AE is naturally associated with the ability of dislocations to move, i.e. with ductility, while the burst type AE is associated with fracture events. Hardened or hydrogen embritteled materials, where dislocation glide is restricted or impeded, are prone to micro- and macro-cracking since the cracks

serve as primary mechanisms of local stress relaxation instead of dislocation motion. Hence, with increased hardening and reduced ductility the continuous AE component, such as that exemplified in Fig. 4a, should decrease while the discrete component, Fig. 4b and c, should increase. Indeed this is observed clearly in the present work. The ductility reduction caused either by hydrogen charging or low strain rate results in the remarkable increase of the burst type AE signals during testing (cf. N in Table 2). Thus, the increase in the intensity accumulation of discrete AE is noticed. To illustrate this trend the dependence of N /max on ε˙ is shown in Fig. 5 in logarithmic coordinates for all types samples. Apparently, the N /max ratio alone does not represent correctly the picture of damage accumulation during straining since the cumulative AE N is not suitable for revealing irregularities in the AE time-series. Fig. 6 demonstrates the kinetics of AE accumulation in steel 70 in dependence of the strain rate. One can notice that the strain rate reduction down to 3 × 10−6 s−1 gives rise to the most rapid AE accumulation which occurs in a steplike manner, e.g. Fig. 6a and c curve 3. Besides, we should underline that for all types specimens – reference, hydrogenated and dehydrogenated – tested at relatively slow strain rates of ε˙ = 3 × 10−4 and/or 3 × 10−6 s−1 the first notable jump of the number of AE signals is observed only after a certain considerable time interval of nearly complete “silence” when AE is virtually absent, Fig. 6 curves 2, 3. As the strain rate decreases, this time interval becomes shorter and finally at ε˙ = 3 × 10−3 s−1 no incubation period is observed, i.e., AE commences almost immediately after the beginning of loading and then the signal accumulation occurs gradually until fracture, Fig. 6a, b, curve 1. The AE count reflects only the fact of AE arrival, i.e. it says only that a certain local structural transformation, e.g. crack, has occurred within the sample, but it does not say anything about the scale of damage. In the contrast, the AE amplitude (or rms voltage, or energy) can be related to the size of the defect and its velocity [31,32]. It is interesting that although, generally speaking, the AE count and energy are statistically independent variables, the kinetics of AE energy accumulation for discrete AE signals is quite similar to that of AE counts and this will be discussed below. The accumulation of AE bursts for hydrogen charged specimens is seen to be gradual and continuous even at the low strain rate ε˙ = 3 × 10−6 s−1 in contrast with the step-like AE accumulation curves for the reference unplated and DHT specimens, cf. curves 3 in Fig. 6. This is straightforwardly explained by the fact that the time interval between jumps in the N curve at high CH is so small that individual AE burst overlap and become indistinguishable, cf. also Fig. 4c illustrating AE time series of discrete bursts following each other with high intensity.

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Fig. 2. Whole thickness (a) fracture surface of unplated specimen tested at high strain rate ε˙ = 3 × 10−3 s−1 . Fragment (b) of the fracture surface zoomed from a selected area A marked on (a) illustrating that crack occurred by ductile decohesion of sulfide inclusions; (c) EDS spectrum taken from one of inclusions showed on (b). The numeric values corresponding to respective concentrations of elements of interest in wt.% and at.% are shown in the inset.

The rms voltage, URMS , reveals a characteristic AE peak in the quasi-elastic part of the loading curve for all specimens at high and middle strain rates of 3 × 10−3 and 3 × 10−4 s−1 , respectively, Figs. 7a–d and 8a, b. However, we should stress that the AE peak for charged specimens is 5–10 times higher than that for the reference and DHT specimens. Furthermore, the position of the peak is shifted to significantly lower loads, cf. also Table 3. When ε˙ reduces from 3 × 10−3 to 3 × 10−4 s−1 the AE peak ∗ height, URMS , reduces for all specimens, which is a common trend in the AE dependence on the strain rate [22,33]. However, this ∗ URMS reduction occurs in charged specimens significantly differently than in other ones: in the reference and DHT specimens the ∗ URMS value decreased by a factor of 10 whereas in the cathodically charged specimens this reduction was by far smaller, i.e. by fac∗ peak for reference tor of 5–7, Table 3. The position of the URMS and DHT specimens with respect to the loading curve does not change substantially. However, the load corresponding to the AE onset increases when the strain rate decreases. This load and load of URMS peak in hydrogen charged specimens is consistently smaller than that in reference and DHT specimens. In all cases the high energy AE bursts, such as those shown in Fig. 4b and c, are seen.

Their “instant” amplitudes or rms values usually exceed by far the average height of the AE rms peak seen at the onset of deformation, Figs. 7c, d and 8b. The further decrease in ε˙ by two orders of magnitude to 3 × 10−6 s−1 results in vanishing of continuous AE for all types samples. The URMS magnitude also reduces for the discrete AE bursts, Figs. 7e, f and 8c. However, the hydrogen charged specimens still exhibit a pronounced AE peak which is created by a large amount of intensive bursts. If compared to higher strain rate ε˙ = 3 × 10−4 s−1 the height of the peak is lower by a factor of 2 or 4, while the position of the peak corresponds to a smaller stress, Fig. 7f, Table 3. 4. Discussion The experiments showed that the hydrogen content increases and the ductility decreases substantially after electroplating of steel 70. The subsequent dehydrogenating heat treatment results in reduction of the hydrogen concentration and this recovers ductility at the high strain rate 3 × 10−3 s−1 . However, the embrittlement becomes more pronounced with strain rate reduction for all specimens. This picture is typical of the so-called reversible hydrogen

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Fig. 3. Whole thickness (a) fracture surface of plated specimen tested at low strain rate ε˙ = 3 × 10−6 s−1 . B-area showed the brittle fracture zone of ellipse form exists in central part of specimen. Fragment (b) of the fracture surface zoomed from a selected area A marked on (a) illustrating part of the crack which occurred normal to fracture plane in brittle intercrystalline manner (c). (For interpretation of the references to color in text, the reader is referred to the web version of the article.)

embrittlement associated with crack initiation and propagation facilitated by diffusion-active hydrogen [14,34]. We have seen that the ductility reduction is accompanied by the increasing area fraction of brittle intergranular fracture surface, Figs. 2 and 3, and, simultaneously, by the increasing amount of discrete type AE, Fig. 4b, cf. Table 2. It is therefore plausible to suppose that these AE transient signals are originated from microscopic brittle fracture processes such as intergranular cracking along the boundaries separating austenitic grains. Apparently, the number of transient events of this kind should scale with the intensity of HAC. The N /max parameter in this respect serves as a measure of microcracking intensity, which increases with the increase in CH and with the decrease in ε˙ according to Fig. 5. It is well known that the crack growth under HAC conditions is facilitated by hydrogen motion toward the crack either by lattice diffusion or by dislocation transport. The crack propagation is controlled by the stress intensity factor KI , which depends on the crack geometry (length and shape) and the average far field stress. The crack advance occurs in a jump-like manner when a local stress intensity factor at the crack tip exceeds a certain critical value Kc . H ahead of This value is lower when the hydrogen concentration Ccr the crack tip is higher. Hence, Kc is strongly affected by a relation H and the average flow stress . When the strain rate between Ccr is low and the hydrogen concentration is high, the conditions for cracking can be fulfilled even at low  values. Upon loading, the flow

H concentration can stress increases rather slowly whereas the Ccr potentially reach very high values during short time intervals due to high mobility of H atoms. Therefore, hydrogen assisted multiple cracking occurs in an avalanche-like manner with high frequency of crack occurrence. This corresponds to a characteristic AE behavior represented by curves 2 and 3 in Figs. 6b and 7d and f. If the hydrogen concentration is not very high, the time which is required for hydrogen delivery to the crack tip increases. Accordingly, the time intervals between the subsequent crack jumps and corresponding AE bursts increase too, which is reflected by the curves 3 in Figs. 6a, c, 7c, e and 8c. At higher strain rates the substantially smaller number of hydrogen atoms has enough time to diffuse to the zone of the maximum tensile stresses around the stress risers such as crack tips, interphase boundaries, grain boundaries, etc. Therefore, brittle fracture occurs only in the regions where the initial local hydrogen concentration is high enough. As was stated above, the behavior of the cumulative AE count and energy for discrete AE events N and energy E is similar. This is justified in Fig. 9 where E is plotted versus N for plated specimens at different strain rate. The proportionality between these two variables is evident; the linear regression fit is obtained by the least square method with the correlation coefficient r2 > 0.98. This makes it reasonable to conclude that the average energy “per event” for hydrogen charged specimens is nearly constant during the test, i.e. despite the gradually increasing load, the scale of the

1166 700 807

1284 984 1300

– 0.08 –

– 256 –

– 631 –

413

1266 590 1307 731 150 725

0.04 0.33 0.03

AE peak height, ∗ ,V URMS AE peak height, ∗ ,V URMS Load corresponding to AE ∗ maximum, PAE , N Load corresponding to AE onset, PAE , N

0.35 2.10 0.25

Fig. 4. Typical AE waveforms in hydrogen charged specimens: (a) low amplitude continuous AE realization, (b) a short time series of overlapping burst-type signals of high amplitude, (c) discrete AE “burst”-type signal.

Reference Hydrogen charged Dehydrogenated

AE peak height, ∗ ,V URMS

3 × 10−3

˙ s−1 Initial strain rate ε, Specimen state

Table 3 Parameters of acoustic emission peak in dependence of hydrogen charging and strain rate.

3 × 10−4

Load corresponding to AE onset, PAE , N

Load corresponding to AE ∗ maximum, PAE , N

3 × 10−6

Load corresponding to AE onset, PAE , N

Load corresponding to AE ∗ maximum, PAE , N

E.D. Merson et al. / Materials Science and Engineering A 550 (2012) 408–417

AE generating events remains constant and variations in the rms voltage (or energy) are attributed to different activity of AE sources in terms of the number of sources activated per unit time. The different slopes of the fitting lines show that the average energy “per event” is strain rate dependent: the energy slightly increases with strain rate; however, this dependence is weak and non-linear. Similar dependences are seen for the other types samples (reference and dehydrogenated) although with slightly smaller regression coefficient.

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Fig. 5. Average intensity of discrete AE accumulation, N /max , during mechanical testing depending on the strain rate for different specimens.

The continuous AE signal is clearly observed at strain rates of ε˙ = 3 × 10−3 and 3 × 10−4 s−1 ; cf. the typical waveform shown in Fig. 4a. This continuous AE forms a pronounced peak in the quasi elastic region as is very commonly observed for a wide range of metals due to plastic (or micro-plastic) deformation [22]. Moreover, it has been well established that the energy of AE caused by dislocation motion during uniform plastic deformation should depend linearly on the √ strain rate [33], i.e. the AE rms voltage should be ˙ This dependence is however not seen in the refproportional to ε. erence and DHT samples: as ε˙ reduces from 3 × 10−3 to 3 × 10−4 s−1 ∗ the height of the AE peak URMS decreases by a factor of 10, cf. Table 3, i.e. the AE energy appears to be roughly proportional to ε˙ 2 and not to ε˙ as is expected for uniform plastic deformation. The conclusion is that the continuous AE, which is observed in the present work, is not associated with uniform plastic deformation. This seems to be quite reasonable for high strength steels. The AE peak vanishes in the reference and DHT specimen at ε˙ = 3 × 10−6 s−1 . In view of the above, the latter is not surprising because with strain rate reduction by two orders of magnitude from 3 × 10−4 to 3 × 10−6 s−1 , the AE rms peak ∗ should decrease by a factor of 100 or so, i.e. the URMS is expected to fall below 4 × 10−4 V that is less than the laboratory noise level in the current experiments (of 10−3 V after amplification). As has been shown above, the microstructure of the present high strength steel is featured by a large volume fraction of MnS inclusion serving as a root cause for microcrack nucleation. Besides, as these sulfides are natural stress risers, plastic deformation tend to localize around and to commence at lower far filed stresses than in the rest of the sample. Since the onset of the AE URMS peak is observed during the quasi elastic loading stage, it is plausible to associate the continuous AE observed during deformation of the steel 70 with plastic deformation localized at inclusions. While the AE maximum is very poorly defined in the reference and DHT specimens at the slow strain rate, the hydrogen charged ˙ comspecimens exhibit a clearly pronounced AE peak at the same ε, pare Fig. 7e and f. However, in the coated samples the AE peak consists of a large number of discrete bursts, which are originated from hydrogen-induced cracking. At higher strain rates the AE peak in hydrogen charged specimens is significantly (5–10 times) higher than in the reference and DHT specimens. This is explained by superposition of the continuous AE caused, presumably, by localized plastic deformation at the stress raisers and the burst-type AE associated with brittle crack advance. This superposition of two processes accounts for the apparent luck in proportionality between the AE peak magnitude and the strain rate in hydrogen charged specimens. The intensity of microcracking increases with strain rate reduction in plated specimens, i.e. the contribution of this process in to resultant AE increases while the continuous AE

Fig. 6. Accumulation of AE transients during three point bending testing of the reference (unplated) specimens (a), hydrogen charged (plated) (b) and dehydrogenated specimens (c) at ε˙ = 3 × 10−3 s−1 (1), ε˙ = 3 × 10−4 s−1 (2) and ε˙ = 3 × 10−6 s−1 (3).

component reasonably decreases with decreasing strain rate, Fig. 7. Hence, we can conclude that two competitive mechanisms operate during deformation of the high carbon steel 70: (i) highly localized plastic deformation and (ii) brittle intercrystalline cracking promoted by the hydrogen influence. Both these mechanisms generate AE with markedly different features. The contribution of each mechanism varies depending on metallographic and loading factors. When the hydrogen-induced embrittlement effect is relatively

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Fig. 7. Typical loading curves and AE RMS voltage for high carbon steel specimens at ε˙ = 3 × 10−3 s−1 – (a) reference (unplated), (b) cathodically charged (Zn-plated); ε˙ = 3 × 10−4 s−1 – (c) reference, (d) cathodically charged; ε˙ = 3 × 10−6 s−1 – (e) reference, (f) cathodically charged.

low (e.g. at high strain rate in reference specimens), the continuous AE associated with plastic deformation is dominant, Figs. 7a and 4c. To the contrast, when the effect of hydrogen is significant (at low strain rate in hydrogenated specimens) AE appears as an intensive flow of high amplitude pulses, Figs. 7f and 4a. In the case when both plastic deformation and brittle fracture take place in approx∗ either increase, Fig. 7b and imately equal proportions the AE URMS d, or one can see a peak of continues AE superimposed with high energy burst signals, Fig. 7c and 8b.

Interestingly that, despite the fact that hydrogen concentration in DHT specimens is sufficiently high mechanical behavior and the accompanying AE in these samples at relatively high strain rates 3 × 10−3 and 3 × 10−4 s−1 , Fig. 8a and b, is very much resembling that in the reference samples, Fig. 7a and c. Indeed, embrittlement in the dehydrogenated samples is not observed at these strain rates, cf. Fig. 5 – the magnitudes of N /max dehydrogenated and reference are pretty much alike. Nevertheless, at low strain rate the ductility of the dehydrogenated samples reduces whereas the

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Fig. 9. Dependence of cumulative AE energy of discrete events, E , on cumulative AE events N for hydrogen charged specimen at low strain rate.

the HAC after dehydrogenation is pronounced only at the slowest strain rate, Fig. 8c, which is accompanied by the increased AE activity. At higher initial strain rates, ε˙ = 3 × 10−3 and 3 × 10−4 s−1 , the only low peak of continuous AE is observed for both reference and dehydrogenated samples. This AE behavior is attributed to localized plastic deformation. It is well known that hydrogen charging occurs non-uniformly during electrolytic plating: hydrogen atoms tend to sit predominantly in the subsurface layer so that the concentration gradients may be steep. Moreover, the sample surface experiences maximum stresses (tensile on the one side of the sample and compressive on the other) during the three-point-bending testing mode. Hence from the very beginning of the test a large portion of hydrogen atoms is located in the zone of maximum tensile stresses. Therefore, HAC may occur almost immediately after the beginning of loading even at the high strain rate. In this case microcracking overlaps with localized plastic deformation which gives rise to additional increase in the URMS maximum and its notable shift to lower loads. As bending proceeds, the fracture process appears to be strain rate dependent because it is being controlled by the hydrogen transport toward developing microcracks. The microcracks nucleate at the specimen surface and then propagate in a brittle intercrystalline manner through the sample cross-section under the influence of hydrogen and increasing stresses, forming thereby the characteristic fracture zone shown schematically in Fig. 3a. The macroscopic shear crack, which is seen in Fig. 3a, forms in the central part of all tested specimens, depending on the hydrogen content and strain rate: ductile fracture along sulfides interphase boundaries at low CH and high ε˙ or brittle intercrystalline fracture ˙ at high CH and low ε. Fig. 8. Typical loading curves and AE RMS voltage for dehydrogenated specimens at ε˙ = 3 × 10−3 s−1 (a); ε˙ = 3 × 10−4 s−1 (b); ε˙ = 3 × 10−6 s−1 (c).

fraction of the burst type AE increases sharply in line with the common trends for the other types samples. Thus, in compliance with the gas analysis the behavior of mechanical and AE characteristics signifies that the dehydrogenation heat treatment does not result in complete removal of hydrogen from the samples. This is reasonable since the Zn coating serves as a diffusion barrier for hydrogen. However, the distribution of hydrogen is expected to be more uniform within the sample volume after dehydrogenation. For notable ductility reduction it is essential to redistribute hydrogen and to deliver it to the zones of maximum tensile stresses which then serve as crack nucleation sites. This process takes time. Therefore,

5. Conclusions 1. Hydrogen charging of the high strength steel 70 exerts a profound effect on the ductility and damage accumulation during three-point-bending testing at a constant loading rate. The acoustic emission accompanying the dynamics of microfracture is found useful for real-time capturing of HAC and for in situ monitoring of damage evolution. 2. The cumulative AE event count correlates strongly with the specimen ductility, which is in turn controlled by the strain rate, amount and distribution of hydrogen, and therefore with the fraction of brittle component on the fracture surface. The average energy “per event” is nearly constant during the test, which is indicative that the scale of the elementary fracture events remains constant upon loading. Variations in the AE rms

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voltage are attributed to different activity of AE sources, depending on the strain rate. 3. Two competitive mechanisms operate during three-point bending deformation of the high carbon steel 70: (i) highly localized plastic deformation and (ii) brittle intercrystalline cracking promoted by hydrogen influence. Contributions of both these mechanisms are strongly dependent on the strain rate, and the amount and distribution of the dissolved hydrogen. Finally, it might be worth noting that since both the HAC phenomenon and the reflecting AE signal have been shown to be strain rate sensitive, there seems to be a possibility to assess the propensity for HAC from relatively high strain rate tests, thereby reducing testing time and cost substantially. The experimental issues which need to be resolved before the AE technology can be used to this extent in practice include increasing the signal-to-noise ratio through optimization of the sensor response, using cuttingedge low noise preamplifiers and advanced signal detection and recognition techniques. As a future scope, the present study can be extended to address these issues. Acknowledgements Financial support from the Russian Ministry of Education and Science through the grant-in-aid No. 11.G34.31.0031 is greatly appreciated. Special thanks go to Dr. I. Yasnikov for his skillful help with SEM observations. References [1] [2] [3] [4]

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