Next Article in Journal
Hydrogen Blending as a Transitional Solution for Decarbonizing the Jordanian Electricity Generation Sector
Previous Article in Journal
Hydrogen-Rich Gaseous Mixture for Enhanced Combustion in a Flex-Fuel Engine: An Experimental Analysis
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
Journals
Hydrogen
Volume 6
Issue 4
10.3390/hydrogen6040100
hydrogen-logo
Submit to this Journal Review for this Journal Propose a Special Issue
► Article Menu

Article Menu

share Share announcement Help format_quote Cite question_answer Discuss in SciProfiles
first_page
settings
Order Article Reprints
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Hydrogen Diffusivity and Hydrogen Traps Behavior of a Tempered and Untempered Martensitic Steel

by
Edgar López-Martínez
1,
Samuel Eduardo Salud-Ordon
2,
Octavio Vázquez-Gómez
3,*,
Miguel Iván Dávila-Pérez
3,
Julio C. Villalobos
3 and
Jesus Israel Barraza-Fierro
4
1
Campus Tehuantepec, Universidad del Istmo, Santo Domingo Tehuantepec 70614, Mexico
2
Centro de Investigación y Desarrollo Tecnológico en Electroquímica, Pedro Escobedo, Querétaro 76703, Mexico
3
Tecnológico Nacional de México/IT Morelia, Morelia 58120, Mexico
4
Escuela Preparatoria, Universidad La Salle Nezahualcóyot, Nezahualcóyot 57205, Mexico
*
Author to whom correspondence should be addressed.
Hydrogen 2025, 6(4), 100; https://doi.org/10.3390/hydrogen6040100
Submission received: 19 September 2025 / Revised: 24 October 2025 / Accepted: 26 October 2025 / Published: 4 November 2025
Browse Figures
Versions Notes

Abstract

The effect of tempering temperature and tempering time on the density of hydrogen traps, hydrogen diffusivity, and microhardness in a vanadium-modified AISI 4140 martensitic steel was determined. Tempering parameters were selected to activate the second, third, and fourth tempering stages. These conditions were intended to promote specific microstructural transformations. Permeability tests were performed using the electrochemical method developed by Devanathan and Stachurski, and microhardness was measured before and after these tests. It was observed that hydrogen diffusivity is inversely proportional to microhardness, while the density of hydrogen traps is directly proportional to microhardness. The lowest hydrogen diffusivity, the highest trap density, and the highest microhardness were obtained in the as-quenched condition and the tempering at 286 °C for 0.25 h. In contrast, tempering at a temperature corresponding to the fourth tempering stage increases hydrogen diffusivity and decreases the density of hydrogen traps and microhardness. However, as the tempering time or temperature increases, the opposite occurs, which is attributed to the formation of alloy carbides. Finally, hydrogen has a softening effect for tempering temperatures corresponding to the fourth tempering stage, tempering times of 0.25 h, and in the as-quenched condition. However, with increasing tempering time, hydrogen has a hardening effect.

1. Introduction

Steel is one of the primary structural materials with which bridges, buildings, ships, trains, cars, surgical materials, prostheses, etc., are built. Steels are widely used in the construction and automotive industries, leading to notable steel development. Tempered martensitic steel has high hardness, toughness, and fatigue resistance. It has been widely used in industries that require high-strength steel. However, in an environment with gaseous or ionic hydrogen, atomic hydrogen can be adsorbed, absorbed, diffused, and trapped inside tempered martensitic steel. These can lead to what is known as hydrogen embrittlement (HE) [1,2]. It is accepted that, with the increase in mechanical properties, the HE susceptibility increases [1,3,4,5,6]; however, the improvement in mechanical properties is also related to the increase in hydrogen traps, which, in turn, is related to a decrease in hydrogen diffusivity [7]. HE susceptibility has been studied in several high-strength steels using hydrogen permeation tests [8,9,10,11]. This method can obtain information such as the effective diffusivity coefficient, the hydrogen concentration, and the density of hydrogen traps [12,13].
One way to decrease HE susceptibility is by forming alloy carbides [14,15,16,17], which function as hydrogen traps. Some studies suggest that transition carbides can also act as hydrogen traps [18,19], helping to decrease hydrogen diffusivity. If traps are irreversible, HE susceptibility can be controlled. However, HE susceptibility can be promoted if they are reversible traps since these traps can act as a hydrogen source [20,21]. Also, when the traps reach saturation, the HE susceptibility is restored [22].
In steels alloyed with Mo, V, Cr, Nb, Ti, and W, transition carbides are formed during low-temperature tempering, and alloy carbides can be formed during higher-temperature tempering. The main objective of tempering is to improve the toughness of steel with the minimum sacrifice of mechanical properties.
In the tempering of alloyed steels, there are four overlapping tempering stages: precipitation of transition carbides (first stage), decomposition of retained austenite (second stage), replacement of transition carbides by cementite and partial loss of martensite tetragonality (third stage), and formation of alloy carbides (fourth stage). Secondary hardening may occur depending on factors such as chemical composition, heating rate, tempering temperature, and tempering time. This phenomenon is associated with the formation of alloy carbides in the fourth tempering stage [23,24,25,26,27]. In addition to increasing the mechanical properties, alloy carbides can decrease hydrogen diffusivity, acting as irreversible traps [28], where the amount and size of the alloy carbides and the martensitic matrix control the hydrogen diffusivity [29].
In applications where steels with high mechanical properties are exposed to hydrogen environments, it is crucial to minimize HE susceptibility. To achieve this, more research is needed to optimize the mechanical properties in the presence of hydrogen, particularly regarding the density of hydrogen traps and hydrogen diffusivity, especially in martensitic steels. Therefore, this study aims to investigate how tempering temperature and tempering time can enhance the density of hydrogen traps and reduce hydrogen diffusivity without compromising the microhardness in V−Mo martensitic steel.

2. Materials and Methods

2.1. Material and Heat Treatment

A commercial alloy AISI 4140 and a vanadium ferroalloy (82.97V–0.37Al–1.01Si–0.18C% by w.) were melted in a VIP 75 induction furnace (Inductotherm, Droitwich, UK) to adjust the chemical composition to 0.30% of V (% by mass). Once the chemical composition was adjusted, the molten steel was poured into ingot molds. The chemical composition (Table 1) of the ingots was measured by optical emission spectrometry (SPECTROMAXx LMM14, Kleve, Germany). The ingots obtained were homogenized with a homogenization heat treatment, austenitizing at a temperature of 1000 °C for 20 min. Precision cuts were made to obtain square membranes with an area per side of approximately 0.483 × 10−3 m2. These membranes were roughed and polished to obtain an approximate thickness of 1.0 × 10−3 m. Subsequently, the membranes were austenitized at 1000 °C for 40 min, and then they were quenched in stirred oil for 5 min (as-quenched condition). Next, the membranes were tempered for 0.25 h at three different temperatures: (1) 286 °C, (2) 530 °C, and (3) 580 °C. Each membrane was introduced into the furnace until it reached the indicated temperature, ensuring each had an effective 0.25-hour tempering treatment. Once the tempering time was completed, the membranes were allowed to cool in still air. Additionally, at a tempering temperature of 530 °C, tempering was carried out at three other different times, as follows: (1) the membranes were introduced in the turned-off furnace; (2) the furnace was turned on, programing the setup at 530 °C; (3) once the desired temperature was reached, one membrane was removed after 0.25 h (0.25 h + Δt, where Δt is the time it took the sample to reach the temperature of 530 °C), another after 2 h (2 h + Δt), and the last one after 6 h (6 h + Δt). For these tests, the desired temperature was reached in approximately 35 min, that is, Δt ≈ 0.58 h. Table 2 shows a summary of the tempering parameters used. The tempering at 286 °C occurs within the second and third tempering stages, while the treatments at 530 °C and 580 °C correspond to the fourth tempering stage.

2.2. Hydrogen Permeation

The quenched and tempered membranes were ground with silicon carbide (SiC) paper with the 240, 320, 400, and 600 grit sequence. Electrochemical hydrogen permeation tests were performed using the standardized double hydrogen permeability cell [30]. The experimental setup consisted of a charging cell containing a 0.5 M sulfuric acid solution with 0.2 g of arsenic trioxide per liter to prevent hydrogen recombination and an oxidation cell with a 0.1 M sodium hydroxide solution. The cell configuration is shown in Figure 1.
The solutions were purged with argon to deaerate the solution for 0.5 h; then, the permeation tests were performed using an electrochemical workstation model CS350 (Wuhan Corrtest Instruments Co., Ltd., Wuhan, China). During the tests, the temperature of the solutions in the charging cell and oxidation cell was kept constant (≈20 °C). The potential set at the control value was +300 mV concerning the calomel-saturated electrode. The magnitude of the current density in the charging cell was set at 40 mA·cm−2. A first series of experiments was performed for each tempering temperature and tempering time condition. To verify reproducibility, a second series of experiments were conducted under the same conditions.

2.3. Microstructural Characterization

The microstructure was revealed in the as-annealed condition and as-quenched condition using standard metallographic preparation. Cutting and roughing was conducted with the sequence of sandpapers 240, 320, 400, and 600, followed by polishing with 1.0 μm particle size alumina, then with 0.3 μm alumina. To reveal the microstructure, an etching with Nital 2 was carried out for 5 s, followed by rinsing with ethyl alcohol. Subsequently, the sample was etched with a 10% by weight sodium bisulfate solution for 5 s and rinsed again with ethyl alcohol. Microstructure images were acquired with the optical microscope Axio Observer (Zeiss Group, Oberkochen, Germany).

2.4. Microhardness

Microhardness tests were conducted on the membranes before and after hydrogen permeation in the as-quenched and as-tempered conditions. These tests were performed using a VLL-101 Microdurometer (Mitutoyo, Kawasaki, Japan) with a force of 300 g and a dwell time of 15 s. A total of 10 measurements were taken for each sample.

3. Results

The microstructure of the steel in the as-annealed condition (Figure 2) consists of proeutectoid ferrite and pearlite grains, along with the presence of precipitates. In contrast, Figure 3 illustrates the microstructure of the steel in the as-quenched condition, which is comprised of martensite.
Table 3 presents the mean microhardness obtained in the as-quenched condition, as well as the mean microhardness measured after each tempering condition (as-tempered condition), denoted as microhardness before hydrogen permeation (MBHP). After the permeation tests, slight corrosion was observed on both sides of the membrane; therefore, the surfaces were carefully cleaned with a microfiber cloth prior to re-measuring the microhardness. These values were referred to as microhardness after hydrogen permeation (MAHP).
In order to determine whether the microhardness values are statistically different or equivalent in the as-quenched condition and for each tested tempering condition and before and after hydrogen permeation, Fisher’s test and Student’s t-test were applied. First, Fisher’s test was performed to assess the homogeneity of variances between groups. Depending on this outcome, Student’s t-test was subsequently applied to compare the mean microhardness, assuming equal or unequal variances as appropriate. A significance level of p < 0.05 was used in both tests.
From this analysis, at 530 °C, statistical testing indicated no significant difference between MBHP and MAHP for tempering times of 0.25 h + Δt and 2 h + Δt, whereas, at all other conditions, the two parameters differed significantly. This indicates that hydrogen does not have a significant effect on microhardness when the tempering conditions are 530 °C and the tempering time is up to 2 h.
The atomic hydrogen permeation flux (J(t)) was determined according to the equation:
J t = I t / A F
where I(t)/A is the current density and F is Faraday’s constant. In Figure 4 and Figure 5, the normalized atomic hydrogen permeation flux (J(t)/Jmax) is observed, where Jmax is the atomic hydrogen permeation flux at steady state. Figure 4 shows J(t)/Jmax for the membrane in the as-quenched condition and the membranes tempered for 0.25 h at three different temperatures (286 °C, 530 °C, and 580 °C). Figure 5 presents J(t)/Jmax for the as-quenched condition and the membranes tempered at 530 °C at three different tempering times (0.25 h + Δt, 2 h + Δt, and 6 h + Δt). In both figures, it can be seen that, for the different tempering temperatures and tempering times, different times are required to reach Jmax. Therefore, it is established that hydrogen has different permeation rates due to the different traps present in the microstructure (grain boundaries, dislocations, residual stresses, and alloy carbides) that occur in the quench and during the different stages of tempering [12,31].
From the analysis of the permeation transients, the hydrogen diffusivity was determined and, based on this information, the density of hydrogen traps was calculated.
Hydrogen diffusivity (effective diffusivity of atomic hydrogen, Deff) was determined using the elapsed time method as a function of tempering temperature and tempering time as follows:
D e f f = L 2 6 t l a g
where tlag is the time to achieve a value of J(t)/Jmax = 0.63 and L is the membrane thickness (1.00 ± 0.02).
Table 4 presents the hydrogen diffusivity for both the as−quenched condition and all the tempering conditions. In the as-quenched condition, a diffusivity of 4.16 × 10−7 cm2·s−1 was measured. In contrast, for the different tempering conditions, diffusivity ranged from 1.15 × 10−6 cm2·s−1 (at 286 °C for 0.25 h) to 9.85 × 10−7 cm2·s−1 (at 530 °C for 6 + Δt). Omura and Oyama [32] determined hydrogen diffusion coefficients for martensitic steels in the as-quenched condition, finding values between 10−6 and 10−7 cm2 s−1, depending on the chemical composition. The authors noted that the addition of Cr and Mo significantly decreases diffusivity, as hydrogen interacts with these elements in solid solution. In addition, Cupertino−Malheiros et al. [33] reported hydrogen diffusion coefficients ranging from 1.0 × 10−6 to 4.4 × 10−7 cm2·s−1 for low-alloy steels tempered within the range of the fourth tempering stage. These values are consistent with those obtained in the current study, further supporting the reliability of the diffusivity values measured under different tempering conditions as well as in the as-quenched condition.
The hydrogen diffusivity, MBHP, and MAHP are presented in Figure 6a for the membrane in the as-quenched condition and the membranes tempered for 0.25 h at three different temperatures (286 °C, 530 °C, and 580 °C) and in Figure 6b for the membrane in the as-quenched condition and the membranes tempered at 530 °C for three different tempering times (0.25 h + Δt, 2 h + Δt, and 6 h + Δt). In both figures, it can be seen that the hydrogen diffusivity is inversely proportional to the microhardness. MBHP decreases as the tempering temperature increases, reaching a minimum at 530 °C compared to the as-quenched condition. However, at a tempering temperature of 580 °C, MBHP starts to increase again (see Figure 6a).
In a previous study, Díaz−Villaseñor et al. [34] reported that, for this steel, the second and third tempering stages overlap, with onset temperatures ranging from 277 °C to 297 °C, whereas the fourth tempering stage begins between 484 °C and 511 °C. Accordingly, tempering at 286 °C promotes microstructural transformations characteristic of the second and third stages. However, the precipitation of transition carbides—typically associated with the first stage of tempering—may also occur under these conditions.
The increase in MBHP at a tempering temperature of 580 °C occurs because the fourth tempering stage is carried out at this temperature [34], which corresponds to the alloy carbide precipitation process and is known as secondary hardening. At the tempering temperature of 530 °C, the fourth tempering stage also occurs, since a local maximum of MBHP (393.3 HV) is observed after 2 h of tempering (2 h + Δt) and an absolute minimum (362.8 HV) at 6 h of tempering (6 h + Δt) (Figure 6b). After hydrogen permeation, with a tempering time of 0.25 h and in the as-quenched condition, the microhardness decreases, except for the membrane tempered at 286 °C, where a slight increase is observed, as seen in Figure 6a. In this same figure, when tempered at 530 °C for 0.25 h, the MAHP is lower than the MBHP. However, as the tempering time increases, the MAHP rises to exceed the MBHP, indicating a hardening effect caused by hydrogen (Figure 6b).
In the as-quenched condition, the hydrogen diffusivity is 4.16 × 10−7 cm2·s−1, the MBHP is 509.6 HV, and the MAHP is 447.6 HV. This indicates that, in the as-quenched condition, hydrogen has a softening effect. For the membranes tempered for 0.25 h, the hydrogen diffusivity has an absolute maximum at a tempering temperature of 530 °C and an absolute minimum at 286 °C. This behavior contrasts with the microhardness measurements, which show a minimum at 530 °C (322.5 HV for MBHP and 284.3 HV for MAHP) and a maximum at 286 °C (481.5 HV for MBHP and 493.6 HV for MAHP). This behavior is related to the microstructure developed in the quenching and tempering processes. For the as-quenched condition, the microstructure is martensite with probably some retained austenite. The latter is known from the microhardness in this condition (MBHP = 509.6 HV). It is known that martensite contains a high concentration of dislocations with a high concentration of residual microstresses, which function as hydrogen traps [28,35], in addition to the fact that the retained austenite also functions as a hydrogen trap [36]. The presence of these traps explains the low hydrogen diffusivity in the as-quenched condition. At 286 °C the hydrogen diffusivity decreases slightly (4.13 × 10−7 cm2·s−1) and the MAHP increases from 447.6 HV (in the as-quenched condition) to 493.6 HV (tempering at 286 °C for 0.25 h). This increase in microhardness is due to the effect of hydrogen because microhardness decreases before hydrogen permeation (Figure 6a). It is known that tempering stages overlap [37,38,39,40] so, at this temperature, the tempering processes corresponding to the first, second, and third stages may be taking place. However, due to the short tempering time, these transformations may not progress significantly, as indicated by slight changes in diffusivity and MBHP. The Student’s t-test comparing the as-quenched condition with tempering at 286 °C showed that the mean microhardness is significantly different (p < 0.05). This finding strongly suggests that microstructural changes occur at this tempering temperature. These changes may include the precipitation of transition carbides, decomposition of retained austenite, loss of martensite tetragonality, and the formation of cementite. Such transformations are characteristic of the first, second, and third tempering stages and contribute to the evolution of microhardness. When tempering at 200 °C, Vieira et al. [37] indicate that the decrease in the microhardness is related to the reduction in solid solution strengthening due to carbon clustering and precipitation of transition carbides.
As the tempering temperature increases to 530 °C with a tempering time of 0.25 h, the hydrogen diffusivity increases and the microhardness decreases (Figure 6a). This temperature is in the region corresponding to the fourth tempering stage, which corresponds to the formation of alloy carbides for alloy steels. As already mentioned, the tempering stages overlap and, due to the short tempering time, it is possible that, at this temperature, the third stage (formation of cementite from transition carbides) and the fourth stage overlap. This can be verified because, at the tempering temperature of 580 °C, there is an increase in MBHP membranes due to the formation of alloy carbides. With the increase in microhardness, the hydrogen diffusivity decreases, and this decrease is known to be related to the formation of alloy carbides since these function as hydrogen traps [41,42,43]. The hydrogen permeation and microhardness of the membranes tempered at 530 °C at different tempering times help verify that alloy carbides are being formed, which modify the hydrogen diffusivity. Figure 6b shows that, as the tempering time increases at a tempering temperature of 530 °C, the MBHP and MAHP initially decrease and then increase up to a tempering time of 2 h + Δt. For the same temperature and a tempering time of 6 h + Δt, the MAHP remains similar to the previous tempering time but is higher than the MBHP, indicating hydrogen hardening.
The density of hydrogen traps (NT) was evaluated using hydrogen diffusivity (Deff) obtained from permeation tests, based on the Oriani-Dong model [44]:
l n D L D e f f − 1 = l n N T N L + E b R T
where DL is the lattice diffusion coefficient of hydrogen, NL is the density of interstitial sites per unit volume, Eb is the hydrogen trap binding energy, R is the universal gas constant, and T is the absolute temperature. Due to the experimental nature of the steel, relevant experimental parameters for NL, DL, and Eb are not available. Furthermore, under tempering conditions, martensite progressively loses its tetragonality, transforming into tempered martensite. Therefore, reference values corresponding to α-Fe were adopted, namely, NL = 7.52 × 1022 sites·cm−3, DL = 1.28 × 10−4 cm2·s−1, and Eb = 0.3 eV [44,45].
Table 4 presents the density of hydrogen traps in the as-quenched condition and as a function of tempering temperature and time. In the as-quenched condition, the density reached 1.96 × 1020 sites·cm−3, reflecting the high dislocation density of martensitic microstructures. After tempering, the density exhibited a dual trend: it could either decrease to 3.06 × 1019 sites·cm−3 due to dislocation recovery or increase up to 2.21 × 1020 sites·cm−3 as a result of alloy carbide precipitation. This contrasting behavior highlights the complex interplay between microstructural recovery and precipitation phenomena in governing hydrogen trapping. Cupertino−Malheiros et al. [33] found that the density of hydrogen traps in low-alloy steel varies with tempering conditions, ranging from 5.9 × 1018 sites·cm−3 to 1.7 × 1020 sites·cm−3. This variation is associated with the precipitation of molybdenum and vanadium alloy carbides. In a chromium−molybdenum steel tempered at 550 °C for 2 h, Zafra et al. [46] calculated the density of hydrogen traps to be 6.73 × 1019 sites·cm−3, with dislocations identified as the primary trapping sites.
Figure 7 shows the density of hydrogen traps, MBHP, and MAHP as a function of tempering temperature and tempering time. Figure 7a shows that most hydrogen traps are in the as-quenched condition and tempered at 286 °C. Concerning the as-quenched condition, the density of hydrogen traps increases at the tempering temperature of 286 °C. At this temperature, transition carbides may be formed, which function as hydrogen traps [47]. With the increase in tempering temperature at 530 °C, there is a decrease in hydrogen traps; later, at 580 °C, there is an increase in hydrogen traps. These two tempering temperatures correspond to the fourth tempering stage, where the precipitation of alloy carbides occurs. Turk et al. [48] indicate that smaller alloy carbides can trap significantly more hydrogen than larger carbides, which has to do with the larger effective area of the precipitate. In Figure 7a, it can be seen that, as the tempering temperature increases, the number of hydrogen traps decreases, reaching a minimum at 530 °C. Although alloy carbides are formed during the tempering process, the decrease in hydrogen traps may be due to the reduction in dislocation density and residual stresses, which act as hydrogen traps [17,49,50]. Subsequently, at 580 °C, the density of hydrogen traps increases due to the more significant presence of alloy carbides. In Figure 7b, it is observed that, as the tempering time increases, the density of hydrogen traps decreases to a minimum at 2 h (2h + Δt) and then increases when the tempering time reaches 6 h (6h + Δt). Once again, the decrease in traps is related to the decrease in dislocations, the decrease in residual stresses and the growth of cementite, and the increase in traps with the more significant presence of alloy carbides. These carbides would be small enough to act as hydrogen traps since hydrogen diffusivity decreases under these conditions (Figure 6b). With this, it is observed that the density of hydrogen traps is directly proportional to the microhardness. Further investigation is required to determine the behavior at tempering times more significant than 6 h.
Hydrogen traps are classified as either reversible or irreversible. Reversible traps have a low trapping energy, allowing hydrogen to diffuse easily, while irreversible traps have a high trapping energy, which makes hydrogen diffusion more difficult. While grain boundaries and dislocations typically act as reversible hydrogen traps, alloy carbides function as irreversible hydrogen traps [43,51,52,53,54]. It is well established that alloy carbides can help reduce the effects of HE [16,55,56]. Among these, vanadium carbides are recognized for having a higher trapping capacity compared to other types of carbides [29,57].
While irreversible traps are often believed to effectively reduce diffusion, this does not always ensure improved resistance to HE. Under specific conditions, such as elevated temperatures, mechanical stresses, or trap saturation, these irreversible traps can become sources of hydrogen and may actually promote embrittlement. For instance, Nagao et al. [22] demonstrated that alloy carbides can help reduce susceptibility to HE. However, once these traps become saturated, the risk of embrittlement returns. This emphasizes the challenge of fully understanding a material’s susceptibility to HE.

4. Conclusions

The effect of tempering temperature and tempering time on the density of hydrogen traps, hydrogen diffusivity, and microhardness in a vanadium-modified AISI 4140 martensitic steel was studied. In the as-quenched condition, the steel presents low hydrogen diffusivity, high density of hydrogen traps, and high microhardness due to the untempered martensite. Depending on the tempering temperature and the tempering time, secondary hardening can occur and decrease the hydrogen diffusivity by increasing the density of hydrogen traps. The following is concluded:
The lowest hydrogen diffusivity occurs in the as-quenched condition and the tempering for 0.25 h at 286 °C, which is related to a high density of hydrogen traps and high microhardness.
At the tempering temperature of 530 °C and 580 °C, the fourth tempering stage is carried out, since, at 580 °C, an increase in microhardness is observed compared to the tempering temperature of 530 °C and, at this last temperature, an increase in microhardness is observed as a function of the tempering time.
For the tempering temperature of 580 °C, hydrogen diffusivity decreases due to the formation of alloying carbides. This can be verified because the microhardness increases concerning the tempering at 530 °C.
Hydrogen softens the untempered martensite and tempered martensite when tempering is performed for a short time (0.25 h), except for the tempering temperature of 286 °C, where a hardening effect occurs.
For tempering at 530 °C, with increasing tempering time, secondary hardening occurs; in addition, hydrogen has a hardening effect.
Microhardness is inversely proportional to hydrogen diffusivity and directly proportional to the density of hydrogen traps.
Finally, if the tempering is conducted at a temperature where the fourth tempering stage is presented, the tempering time may be optimized to maximize microhardness and minimize hydrogen diffusivity by increasing the density of hydrogen traps.

Author Contributions

Conceptualization, S.E.S.-O., E.L.-M. and O.V.-G.; methodology, S.E.S.-O., E.L.-M., O.V.-G., M.I.D.-P., J.C.V. and J.I.B.-F.; software, M.I.D.-P. and J.C.V.; validation, J.I.B.-F.; formal analysis, S.E.S.-O., E.L.-M., O.V.-G., M.I.D.-P., J.C.V. and J.I.B.-F.; investigation, S.E.S.-O., E.L.-M., O.V.-G., M.I.D.-P., J.C.V. and J.I.B.-F.; resources, O.V.-G.; writing—original draft preparation, S.E.S.-O. and E.L.-M.; writing—review and editing, S.E.S.-O., E.L.-M. and O.V.-G.; supervision, O.V.-G., M.I.D.-P. and J.C.V.; project administration, E.L.-M. and O.V.-G.; funding acquisition, O.V.-G. All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

Abbreviations

The following abbreviations are used in this manuscript:
MAHPMicrohardness after hydrogen permeation
MBHPMicrohardness before hydrogen permeation
HEHydrogen embrittlement

References

  1. Venezuela, J.; Liu, Q.; Zhang, M.; Zhou, Q.; Atrens, A. A review of hydrogen embrittlement of martensitic advanced high-strength steels. Corros. Rev. 2016, 34, 153–186. [Google Scholar] [CrossRef]
  2. Lee, S.J.; Ronevich, J.A.; Krauss, G.; Matlock, D.K. Hydrogen Embrittlement of Hardened Low-carbon Sheet Steel. ISIJ Int. 2010, 50, 294–301. [Google Scholar] [CrossRef]
  3. Shi, X.B.; Yan, W.; Wang, W.; Zhao, L.Y.; Shan, Y.Y.; Yang, K. HIC and SSC Behavior of High-Strength Pipeline Steels. Acta Metall. Sin. Engl. Lett. 2015, 28, 799–808. [Google Scholar] [CrossRef]
  4. Adasooriya, N.D.; Tucho, W.M.; Holm, E.; Årthun, T.; Hansen, V.; Solheim, K.G.; Hemmingsen, T. Effect of hydrogen on mechanical properties and fracture of martensitic carbon steel under quenched and tempered conditions. Mater. Sci. Eng. A 2021, 803, 140495. [Google Scholar] [CrossRef]
  5. Venezuela, J.; Hill, T.; Zhou, Q.; Li, H.; Shi, Z.; Dong, F.; Knibbe, R.; Zhang, M.; Dargusch, M.S.; Atrens, A. Hydrogen-induced fast fracture in notched 1500 and 1700 MPa class automotive martensitic advanced high-strength steel. Corros. Sci. 2021, 188, 109550. [Google Scholar] [CrossRef]
  6. Truschner, M.; Pengg, J.; Loder, B.; Köberl, H.; Gruber, P.; Moshtaghi, M.; Mori, G. Hydrogen resistance and trapping behaviour of a cold-drawn ferritic-pearlitic steel wire. Int. J. Mater. Res. 2023, 114, 439–452. [Google Scholar] [CrossRef]
  7. Venezuela, J.; Zhou, Q.; Liu, Q.; Zhang, M.; Atrens, A. Hydrogen Trapping in Some Automotive Martensitic Advanced High-Strength Steels. Adv. Eng. Mater. 2018, 20, 1700468. [Google Scholar] [CrossRef]
  8. Zhang, S.; Wan, J.; Zhao, Q.; Liu, J.; Huang, F.; Huang, Y.; Li, X. Dual role of nanosized NbC precipitates in hydrogen embrittlement susceptibility of lath martensitic steel. Corros. Sci. 2020, 164, 108345. [Google Scholar] [CrossRef]
  9. Rudomilova, D.; Prošek, T.; Luckeneder, G. Techniques for investigation of hydrogen embrittlement of advanced high strength steels. Corros. Rev. 2018, 36, 413–434. [Google Scholar] [CrossRef]
  10. Venezuela, J.; Lim, F.Y.; Liu, L.; James, S.; Zhou, Q.; Knibbe, R.; Zhang, M.; Li, H.; Dong, F.; Dargusch, M.S.; et al. Hydrogen embrittlement of an automotive 1700 MPa martensitic advanced high-strength steel. Corros. Sci. 2020, 171, 108726. [Google Scholar] [CrossRef]
  11. Hu, Y.; Dong, C.; Luo, H.; Xiao, K.; Zhong, P.; Li, X. Study on the Hydrogen Embrittlement of Aermet100 Using Hydrogen Permeation and SSRT Techniques. Metall. Mater. Trans. A 2017, 48, 4046–4057. [Google Scholar] [CrossRef]
  12. Liu, Q.; Venezuela, J.; Zhang, M.; Zhou, Q.; Atrens, A. Hydrogen trapping in some advanced high strength steels. Corros. Sci. 2016, 111, 770–785. [Google Scholar] [CrossRef]
  13. López-Martínez, E.; Vergara-Hernández, H.J.; Flores, O.; Campillo, B. Hydrogen Diffusivity in the Welding Zone of Two High Strength Experimental Microalloyed Steels. ISIJ Int. 2015, 55, 2435–2442. [Google Scholar] [CrossRef]
  14. Lee, J.; Lee, T.; Mun, D.J.; Bae, C.M.; Lee, C.S. Comparative study on the effects of Cr, V, and Mo carbides for hydrogen-embrittlement resistance of tempered martensitic steel. Sci. Rep. 2019, 9, 5219. [Google Scholar] [CrossRef]
  15. Lee, J.; Lee, T.; Kwon, Y.J.; Mun, D.J.; Yoo, J.Y.; Lee, C.S. Role of Mo/V carbides in hydrogen embrittlement of tempered martensitic steel. Corros. Rev. 2015, 33, 433–441. [Google Scholar] [CrossRef]
  16. Lee, J.; Lee, T.; Kwon, Y.J.; Mun, D.J.; Yoo, J.Y.; Lee, C.S. Effects of vanadium carbides on hydrogen embrittlement of tempered martensitic steel. Met. Mater. Int. 2016, 22, 364–372. [Google Scholar] [CrossRef]
  17. Chen, Y.S.; Lu, H.; Liang, J.; Rosenthal, A.; Liu, H.; Sneddon, G.; McCarroll, I.; Zhao, Z.; Li, W.; Guo, A.; et al. Observation of hydrogen trapping at dislocations, grain boundaries, and precipitates. Science 2020, 367, 171–175. [Google Scholar] [CrossRef]
  18. Takahashi, J.; Kawakami, K.; Sakiyama, Y.; Ohmura, T. Atomic-scale observation of hydrogen trap sites in bainite–austenite dual-phase steel by APT. Mater. Charact. 2021, 178, 111282. [Google Scholar] [CrossRef]
  19. Zhu, X.; Li, W.; Hsu, T.Y.; Zhou, S.; Wang, L.; Jin, X. Improved resistance to hydrogen embrittlement in a high-strength steel by quenching–partitioning–tempering treatment. Scr. Mater. 2015, 97, 21–24. [Google Scholar] [CrossRef]
  20. Dadfarnia, M.; Sofronis, P.; Neeraj, T. Hydrogen interaction with multiple traps: Can it be used to mitigate embrittlement? Int. J. Hydrogen Energy 2011, 36, 10141–10148. [Google Scholar] [CrossRef]
  21. Bhadeshia, H.K.D.H. Prevention of hydrogen embrittlement in steels. ISIJ Int. 2016, 56, 24–36. [Google Scholar] [CrossRef]
  22. Nagao, A.; Martin, M.L.; Dadfarnia, M.; Sofronis, P.; Robertson, I.M. The effect of nanosized (Ti,Mo)C precipitates on hydrogen embrittlement of tempered lath martensitic steel. Acta Mater. 2014, 74, 244. [Google Scholar] [CrossRef]
  23. Sackl, S.; Leitner, H.; Clemens, H.; Primig, S. On the evolution of secondary hardening carbides during continuous versus isothermal heat treatment of high speed steel HS 6-5-2. Mater. Charact. 2016, 120, 323–330. [Google Scholar] [CrossRef]
  24. Liu, B.; Qin, T.; Xu, W.; Jia, C.; Wu, Q.; Chen, M.; Liu, Z. Effect of tempering conditions on secondary hardening of carbides and retained austenite in spray-formed M42 high-speed steel. Materials 2019, 12, 3714. [Google Scholar] [CrossRef] [PubMed]
  25. Bała, P.; Pacyna, J. The kinetics of phase transformations during tempering in high-speed steels. J. Achiev. Mater. Manuf. Eng. 2007, 23, 15–18. [Google Scholar]
  26. Claesson, E.; Magnusson, H.; Kohlbrecher, J.; Thuvander, M.; Hedström, P. Evolution of iron carbides during tempering of low-alloy tool steel studied with polarized small angle neutron scattering, electron microscopy and atom probe. Mater. Charact. 2022, 194, 112464. [Google Scholar] [CrossRef]
  27. Hunkel, M.; Dong, J.; Epp, J.; Kaiser, D.; Dietrich, S.; Schulze, V.; Rajaei, A.; Hallstedt, B.; Broeckmann, C. Comparative study of the tempering behavior of different martensitic steels by means of in-situ diffractometry and dilatometry. Materials 2020, 13, 5058. [Google Scholar] [CrossRef]
  28. Shi, R.; Chen, L.; Wang, Z.; Yang, X.S.; Qiao, L.; Pang, X. Quantitative investigation on deep hydrogen trapping in tempered martensitic steel. J. Alloys Compd. 2021, 854, 157218. [Google Scholar] [CrossRef]
  29. Van den Eeckhout, E.; Depover, T.; Verbeken, K. The Effect of Microstructural Characteristics on the Hydrogen Permeation Transient in Quenched and Tempered Martensitic Alloys. Metals 2018, 8, 779. [Google Scholar] [CrossRef]
  30. Devanathan, M.A.V.; Stachurski, Z. The adsorption and diffusion of electrolytic hydrogen in palladium. Proc. R. Soc. Lond. Ser. A 1962, 270, 90–102. [Google Scholar]
  31. Miranda, M.; Campillo, B.; Villalobos, J.C.; Colín, J.; Islas, A.T.; Molina, A.; Serna, S. Analysis of Electrochemical Hydrogen Permeation through AISI Stainless and Welding Zones for Hydrogen Service. Int. J. Electrochem. Sci. 2016, 11, 9340–9354. [Google Scholar] [CrossRef]
  32. Omura, T.; Oyama, T. Effects of the Addition of Alloying Elements on Hydrogen Diffusion and Hydrogen Embrittlement in Martensitic Steel. ISIJ Int. 2024, 64, 620–629. [Google Scholar] [CrossRef]
  33. Cupertino-Malheiros, L.; Oudriss, A.; Thébault, F.; Piette, M.; Feaugas, X. Hydrogen Diffusion and Trapping in Low-Alloy Tempered Martensitic Steels. Metall. Mater. Trans. A 2023, 54, 1159–1173. [Google Scholar] [CrossRef]
  34. Díaz-Villaseñor, P.G.; Vázquez-Gómez, O.; Garnica-González, P.; Vergara-Hernández, H.J.; Campillo-Illanes, B.; López-Martínez, E. Secondary hardening in an Cr-Mo-V steel: Effect of thermal cycle tempering. Mater. Sci. Technol. 2024, 40, 1082. [Google Scholar] [CrossRef]
  35. Luppo, M.I.; Ovejero-Garcia, J. The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel. Corros. Sci. 1991, 32, 1125–1136. [Google Scholar] [CrossRef]
  36. Chan, S.L.I.; Lee, H.L.; Yang, J.R. Effect of retained austenite on the hydrogen content and effective diffusivity of martensitic structure. Metall. Trans. A 1991, 22, 2579–2586. [Google Scholar] [CrossRef]
  37. Vieira, I.; Klemm-Toole, J.; Buchner, E.; Williamson, D.L.; Findley, K.O.; De Moor, E. A Dilatometric Study of Tempering Complemented by Mössbauer Spectroscopy and other Characterization Techniques. Sci. Rep. 2017, 7, 17337. [Google Scholar] [CrossRef] [PubMed]
  38. Primig, S.; Leitner, H. Separation of overlapping retained austenite decomposition and cementite precipitation reactions during tempering of martensitic steel by means of thermal analysis. Thermochim. Acta 2011, 526, 111–117. [Google Scholar] [CrossRef]
  39. Sahoo, G.; Singh, K.K.; Kumar, V. Quenched and tempered high strength steel: A review. J. Met. Mater. Miner. 2020, 30, 19–29. [Google Scholar] [CrossRef]
  40. Morales, E.V.; Vega-Leiva, J.A.; Salinas, H.L.L.; Bott, I.S. Analysis of precipitation mechanisms in tempering of low-alloy steels using the kinetic theory of the overlapping transformations. Phase Transit. 2011, 84, 179–191. [Google Scholar] [CrossRef]
  41. Li, L.; Song, B.; Cheng, J.; Yang, Z.; Cai, Z. Effects of vanadium precipitates on hydrogen trapping efficiency and hydrogen induced cracking resistance in X80 pipeline steel. Int. J. Hydrogen Energy 2018, 43, 17353–17363. [Google Scholar] [CrossRef]
  42. Peral, L.B.; Amghouz, Z.; Colombo, C.; Fernández-Pariente, I. Evaluation of hydrogen trapping and diffusion in two cold worked CrMo(V) steel grades by means of the electrochemical hydrogen permeation technique. Theor. Appl. Fract. Mech. 2020, 110, 102771. [Google Scholar] [CrossRef]
  43. Peral, L.B.; Fernández-Pariente, I.; Colombo, C.; Rodríguez, C.; Belzunce, J. The Positive Role of Nanometric Molybdenum-Vanadium Carbides in Mitigating Hydrogen Embrittlement in Structural Steels. Materials 2021, 14, 7269. [Google Scholar] [CrossRef] [PubMed]
  44. Dong, C.F.; Liu, Z.Y.; Li, X.G.; Cheng, Y.F. Effects of hydrogen-charging on the susceptibility of X100 pipeline steel to hydrogen-induced cracking. Int. J. Hydrogen Energy 2009, 34, 9879–9884. [Google Scholar] [CrossRef]
  45. Villalobos, J.C.; Del-Pozo, A.; Mayen, J.; Serna, S.; Campillo, B. Hydrogen embrittlement suscetibility on X-120 microalloyed steel as function of tempering temperature. Int. J. Hydrogen Energy 2020, 45, 9137. [Google Scholar] [CrossRef]
  46. Zafra, A.; Peral, L.B.; Belzunce, J. Hydrogen diffusion and trapping in A 42CrMo4 quenched and tempered steel: Influence of tempering temperature. Int. J. Hydrogen Energy 2020, 45, 31225–31242. [Google Scholar] [CrossRef]
  47. Shimotomai, M. Heuristic Design of Advanced Martensitic Steels That Are Highly Resistant to Hydrogen Embrittlement by ε-Carbide. Metals 2021, 11, 370. [Google Scholar] [CrossRef]
  48. Turk, A.; San Martín, D.; Rivera-Díaz-del-Castillo, P.E.J.; Galindo-Nava, E.I. Correlation between vanadium carbide size and hydrogen trapping in ferritic steel. Scr. Mater. 2018, 152, 112–116. [Google Scholar] [CrossRef]
  49. Chen, W.; Zhao, W.; Gao, P.; Li, F.; Kuang, S.; Zou, Y.; Zhao, Z. Interaction between dislocations, precipitates and hydrogen atoms in a 2000 MPa grade hot-stamped steel. J. Mater. Res. Technol. 2022, 18, 4353–4366. [Google Scholar] [CrossRef]
  50. Hayato, N.; Tomohiko, H.; Saya, A.; Yuki, S.; Motomichi, K.; Hiroyuki, S.; Ayumi, S.; Ryo, Y.; Takahisa, S.; Eiji, A. Effects of Residual Stress on Hydrogen Embrittlement of a Stretch-Formed Tempered Martensitic Steel Sheet. ISIJ Int. 2021, 61, 1170–1178. [Google Scholar] [CrossRef]
  51. Javeria, U.; Kim, S.J. Investigation of hydrogen embrittlement in steel alloys: Mechanism, factors, advanced methods and materials, applications, challenges, and future directions: A review. J. Mater. Res. Technol. 2025, 38, 1276. [Google Scholar] [CrossRef]
  52. Oliveira, M.S.; Carrasco, J.A.P.; Vilar, E.O.; Duarte, R.C.O.; dos Santos, M.A.; Silva, A.A.; Carvalho, L.H. Heat Treatment Effect on Hydrogen Permeation and Trapping in the API 5CT P110 Steel. Mater. Res. 2023, 26, e20230332. [Google Scholar] [CrossRef]
  53. Cheng, X.Y.; Li, H.; Cheng, X.B. Carbides and possible hydrogen irreversible trapping sites in ultrahigh strength round steel. Micron 2017, 103, 22. [Google Scholar] [CrossRef] [PubMed]
  54. Faye, O.; Szpunar, J.A. Effect of Metal Carbides on Hydrogen Embrittlement: A Density Functional Theory Study. Hydrogen 2024, 5, 137. [Google Scholar] [CrossRef]
  55. Park, J.S.; Jung, S.P.; Kim, S.J. Mitigating hydrogen embrittlement of advanced high-strength steel by controlling carbides (cementite and alloy carbides) and microstructural modification. J. Mater. Res. Technol. 2024, 28, 4703. [Google Scholar] [CrossRef]
  56. Depover, T.; Verbeken, K. Evaluation of the effect of V4C3 precipitates on the hydrogen induced mechanical degradation in Fe-C-V alloys. Mater. Sci. Eng. A 2016, 675, 299. [Google Scholar] [CrossRef]
  57. Seo, H.J.; Heo, Y.-U.; Kim, J.N.; Lee, J.; Choi, S.; Lee, C.S. Effect of V/Mo ratio on the evolution of carbide precipitates and hydrogen embrittlement of tempered martensitic steel. Corros. Sci. 2020, 176, 108929. [Google Scholar] [CrossRef]
Hydrogen 06 00100 g001
Figure 1. Diagram of the permeation cell, consisting of a charging cell and a hydrogen detection cell (oxidation cell). WE: working electrode (membrane); AE: auxiliar electrode; RE: reference electrode (calomel saturated electrode).
Figure 1. Diagram of the permeation cell, consisting of a charging cell and a hydrogen detection cell (oxidation cell). WE: working electrode (membrane); AE: auxiliar electrode; RE: reference electrode (calomel saturated electrode).
Hydrogen 06 00100 g001
Hydrogen 06 00100 g002
Figure 2. Microstructure of steel in the as-annealed condition.
Figure 2. Microstructure of steel in the as-annealed condition.
Hydrogen 06 00100 g002
Hydrogen 06 00100 g003
Figure 3. Microstructure of steel in the as-quenched condition.
Figure 3. Microstructure of steel in the as-quenched condition.
Hydrogen 06 00100 g003
Hydrogen 06 00100 g004
Figure 4. Normalized atomic hydrogen permeation flux for the membrane in the as-quenched condition and for the membranes tempered for 0.25 h at three different tempering temperatures.
Figure 4. Normalized atomic hydrogen permeation flux for the membrane in the as-quenched condition and for the membranes tempered for 0.25 h at three different tempering temperatures.
Hydrogen 06 00100 g004
Hydrogen 06 00100 g005
Figure 5. Normalized atomic hydrogen permeation flux for the membrane in the as-quenched condition and for the membranes tempered at 530 °C for three different tempering times (t + Δt). Δt ≈ 0.58 h.
Figure 5. Normalized atomic hydrogen permeation flux for the membrane in the as-quenched condition and for the membranes tempered at 530 °C for three different tempering times (t + Δt). Δt ≈ 0.58 h.
Hydrogen 06 00100 g005
Hydrogen 06 00100 g006
Figure 6. Hydrogen diffusivity and microhardness for the membrane in the as-quenched condition and (a) for the membranes tempered for 0.25 h at three different temperatures; (b) for the membranes tempered at 530 °C for three different tempering times.
Figure 6. Hydrogen diffusivity and microhardness for the membrane in the as-quenched condition and (a) for the membranes tempered for 0.25 h at three different temperatures; (b) for the membranes tempered at 530 °C for three different tempering times.
Hydrogen 06 00100 g006
Hydrogen 06 00100 g007
Figure 7. Density of hydrogen traps and microhardness for the membrane in the as-quenched condition and (a) for the membranes tempered for 0.25 h at three different temperatures; (b) for the membranes tempered at 530 °C for three different tempering times.
Figure 7. Density of hydrogen traps and microhardness for the membrane in the as-quenched condition and (a) for the membranes tempered for 0.25 h at three different temperatures; (b) for the membranes tempered at 530 °C for three different tempering times.
Hydrogen 06 00100 g007
Table 1. Chemical composition of steel (wt% in mass).
Table 1. Chemical composition of steel (wt% in mass).
%C%Mn%Si%Mo%Cr%V%Ni%Fe
0.3780.7190.2360.1910.9400.2920.071Balance
Table 2. Tempering conditions. (Δt ≈ 0.58 h).
Table 2. Tempering conditions. (Δt ≈ 0.58 h).
Tempering Time (h)Tempering Temperature (°C)
286
0.25530
580
0.25 + Δt530
2 + Δt530
6 + Δt530
Table 3. MBHP and MAHP as a function of tempering time and tempering temperature. Δt ≈ 0.58 h.
Table 3. MBHP and MAHP as a function of tempering time and tempering temperature. Δt ≈ 0.58 h.
Tempering Time
(h)		Tempering Temperature (°C)Mean MBHP
(HV 0.3/15)		Mean MAHP
(HV 0.3/15)
0As-quenched509.6 ± 1.6447.6 ± 5.5
0.25286481.5 ± 3.1493.6 ± 2.2
530322.5 ± 2.8284.3 ± 3.0
580400.8 ± 4.1390.3 ± 2.2
0.25 + Δt530374.4 ± 2.5378.0 ± 4.5
2 + Δt530393.3 ± 3.1401.9 ± 2.7
6 + Δt530362.8 ± 5.6398.6 ± 3.4
Table 4. Diffusivity and density of hydrogen traps as a function of tempering time and tempering temperature. Δt ≈ 0.58 h.
Table 4. Diffusivity and density of hydrogen traps as a function of tempering time and tempering temperature. Δt ≈ 0.58 h.
Tempering Time
(h)		Tempering Temperature (°C)Mean Diffusivity (cm2·s−1)Mean Density of Hydrogen Traps (sites·cm−3)
0As-quenched4.16 × 10−71.96 × 1020
0.252864.13 × 10−72.21 × 1020
5302.71 × 10−63.06 × 1019
5809.85 × 10−78.90 × 1019
0.25 + Δt5301.15 × 10−68.15 × 1019
2 + Δt5301.29 × 10−67.38 × 1019
6 + Δt5305.49 × 10−71.51 × 1020
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

López-Martínez, E.; Salud-Ordon, S.E.; Vázquez-Gómez, O.; Dávila-Pérez, M.I.; Villalobos, J.C.; Barraza-Fierro, J.I. Hydrogen Diffusivity and Hydrogen Traps Behavior of a Tempered and Untempered Martensitic Steel. Hydrogen 2025, 6, 100. https://doi.org/10.3390/hydrogen6040100

AMA Style

López-Martínez E, Salud-Ordon SE, Vázquez-Gómez O, Dávila-Pérez MI, Villalobos JC, Barraza-Fierro JI. Hydrogen Diffusivity and Hydrogen Traps Behavior of a Tempered and Untempered Martensitic Steel. Hydrogen. 2025; 6(4):100. https://doi.org/10.3390/hydrogen6040100

Chicago/Turabian Style

López-Martínez, Edgar, Samuel Eduardo Salud-Ordon, Octavio Vázquez-Gómez, Miguel Iván Dávila-Pérez, Julio C. Villalobos, and Jesus Israel Barraza-Fierro. 2025. "Hydrogen Diffusivity and Hydrogen Traps Behavior of a Tempered and Untempered Martensitic Steel" Hydrogen 6, no. 4: 100. https://doi.org/10.3390/hydrogen6040100

APA Style

López-Martínez, E., Salud-Ordon, S. E., Vázquez-Gómez, O., Dávila-Pérez, M. I., Villalobos, J. C., & Barraza-Fierro, J. I. (2025). Hydrogen Diffusivity and Hydrogen Traps Behavior of a Tempered and Untempered Martensitic Steel. Hydrogen, 6(4), 100. https://doi.org/10.3390/hydrogen6040100

Article Metrics

Back to TopTop