Hydrogen in Steel: Diffusion, Trapping, and Embrittlement Mechanisms

Hydrogen is the smallest and most mobile atom capable of residing interstitially in a steel lattice, yet even sub-ppm concentrations can precipitate catastrophic brittle fracture in otherwise ductile high-strength steels. This article provides a graduate-level treatment of how hydrogen enters steel, how it moves through the BCC and FCC lattice, how trap sites retard or immobilise it, and by which atomistic mechanisms it degrades mechanical properties — culminating in practical guidance on testing, mitigation, and standards compliance.

Hydrogen Sources and Entry into Steel

Hydrogen enters steel as atomic hydrogen (H_ads), not as molecular H₂. Because H₂ cannot dissociate and enter the lattice under normal conditions, the relevant absorption step is always a surface electrochemical or thermal reaction that produces nascent atomic hydrogen before recombination to the gas phase can occur.

Electrochemical entry — corrosion and cathodic reactions

In aqueous acidic or neutral environments, the hydrogen evolution reaction (HER) proceeds in two steps. The Volmer step reduces a proton to an adsorbed hydrogen atom at the metal surface (H⁺ + e⁻ → H_ads). The Heyrovsky or Tafel step then recombines two adsorbed atoms to evolve molecular hydrogen. A fraction of H_ads is absorbed into the lattice before recombination — this sub-surface entry flux is the damaging component. Cathodic overpotential, whether from an external cathodic protection circuit or from coupling to a more noble metal, increases the surface flux of H_ads and therefore the entry rate. Cathodic overprotection below −0.95 V vs Ag/AgCl (commonly specified in DNVGL-RP-F112 for duplex stainless steels) is a recognised HISC risk factor in subsea pipework.

Recombination poisons — the sour service amplifier

In wet H₂S service (“sour service” per NACE MR0175 / ISO 15156), the sulphide ion and H₂S molecule adsorb strongly onto the steel surface and block the recombination step. This forces a greater fraction of H_ads to enter the lattice rather than desorb as H₂. The result is a bulk hydrogen activity orders of magnitude higher than would be predicted from pH alone. Arsenic, antimony, phosphorus, and cyanide ions act similarly. This explains why even moderate partial pressures of H₂S (above 0.3 kPa by NACE definition) require use of sour-service-qualified materials and strict hardness limits (≤250 HV or ≤22 HRC weld HAZ per ISO 15156-2).

Welding and thermal processes

During arc welding, moisture in the electrode coating, flux, or shielding gas dissociates in the high-temperature arc plasma: H₂O → 2H + ½O₂. Atomic hydrogen is highly soluble in liquid iron (at 1600°C, solubility is approximately 27 mL/100g). Upon weld solidification, hydrogen solubility drops precipitously — from 27 mL/100g in liquid to roughly 8 mL/100g in δ-ferrite to less than 1 mL/100g in α-ferrite at room temperature. This supersaturation drives hydrogen toward defects and microstructural heterogeneities, creating the conditions for cold cracking (hydrogen-assisted cracking, HAC) if a susceptible microstructure, residual tensile stress, and sufficient hydrogen concentration coincide — the classical Granjon triangle.

Other thermal hydrogen sources include: bright annealing in hydrogen-rich atmospheres (without adequate post-purge), case hardening by plasma nitriding where hydrogen is a carrier gas, and hydrogen gas pressure testing.

Lattice Diffusion — Interstitial Mechanism and Fick’s Laws

Hydrogen occupies interstitial sites in the iron lattice. In BCC α-ferrite and martensite, it resides preferentially in tetrahedral interstices (T-sites), where the ionic radii mismatch is less severe. In FCC austenite, it occupies larger octahedral interstices (O-sites). The jump frequency between adjacent interstitial sites follows an Arrhenius relationship with activation energy Q_d:

D = D₀ × exp(−Qₓ / RT)

BCC ferrite (25°C):  Dₓ ≈ 1.8 × 10−¹² m²/s,  Qₓ ≈ 7.2 kJ/mol
FCC austenite (25°C): D₀ ≈ 2.0 × 10−³ m²/s,  Qₓ ≈ 48 kJ/mol

Resulting D(25°C):
  Ferrite:   ≈ 1–2 × 10−¹² m²/s  (fast)
  Austenite: ≈ 1–5 × 10−¹⁷ m²/s  (very slow, ~4–5 orders of magnitude lower)

The consequence of this large diffusivity contrast is that hydrogen can traverse a 25 mm ferritic steel section in the order of hours, while the same geometry in austenitic steel retains hydrogen essentially indefinitely under ambient conditions. This is exploited industrially: austenitic stainless steel and nickel alloy cladding over carbon steel pressure vessels functions as a hydrogen permeation barrier, not merely a corrosion barrier.

Fick’s laws applied to steel components

For one-dimensional diffusion under a surface concentration C₀ and zero initial bulk concentration (semi-infinite body approximation), the hydrogen concentration profile is:

C(x,t) = C₀ × erfc( x / (2 √(D×t)) )

where:
  x  = depth from surface (m)
  t  = exposure time (s)
  D  = effective diffusion coefficient (m²/s)
  erfc = complementary error function

In practice, the effective diffusion coefficient D_eff is always lower than the lattice diffusion coefficient D_L because hydrogen is retarded by trap sites. The relationship between D_eff, D_L, and trap density is given by:

D✛?? = Dⁿ / (1 + Nₜ / Nⁿ × exp(Eᴰ / RT))

where:
  Nₜ = trap site density (m−³)
  Nⁿ = lattice interstitial site density (m−³, ≈ 5.1 × 10²¸ for BCC Fe)
  Eᴰ = trap binding energy (J/mol)
  R  = 8.314 J/(mol·K)
  T  = absolute temperature (K)

At room temperature, even a modest trap density of 10²³ m⁻³ with a binding energy of 30 kJ/mol reduces D_eff by roughly one order of magnitude relative to D_L. This is why cold-worked or highly dislocated steels show apparently slower permeation transients in electrochemical permeation experiments (Devanathan–Stachurski cell, ASTM G148).

Hydrogen Trap Sites — Classification and Binding Energies

Not all hydrogen in steel is equally mobile. A fraction resides at trap sites — microstructural features where the local energy is lower than the mean lattice interstitial energy. The depth of this energy well, the trap binding energy E_B, determines whether hydrogen can escape thermally at ambient temperature.

Oriani equilibrium model

Oriani (1970) proposed that hydrogen at trap sites and hydrogen in normal lattice sites exist in local thermodynamic equilibrium. The fractional occupancy of trap sites θ_T is:

θₜ / (1 − θₜ) = θⁿ / (1 − θⁿ) × exp(Eᴰ / RT)

where:
  θₜ = Cₜ / Nₜ  (fractional trap occupancy)
  θⁿ = Cⁿ / Nⁿ  (fractional lattice occupancy)
  Eᴰ     = trap binding energy (J/mol, positive for a trap)

At low hydrogen contents (dilute limit, θ_T << 1 and θ_L << 1) this simplifies to C_T = (N_T/N_L) × C_L × exp(E_B/RT). The model demonstrates that strong traps are essentially saturated well before reversible traps fill, and that they remain occupied even as bulk lattice hydrogen changes.

Reversible (weak) traps

Reversible traps have binding energies in the range 10–60 kJ/mol. Hydrogen can detrap at or near room temperature within the timescale of service or test. Key reversible trap types in steel are:

Irreversible (strong) traps

Irreversible traps have binding energies above approximately 60 kJ/mol. Hydrogen absorbed at these sites at ambient temperature cannot escape without elevated-temperature bake-out. Their role in embrittlement is complex: they immobilise hydrogen and thus reduce the diffusible (mobile) fraction, which may be beneficial in isolation, but under damage accumulation conditions (sustained load, fatigue cycling) they can become sites of hydrogen pressure build-up or decohesion initiation.

Embrittlement Mechanisms

HEDE — Hydrogen-Enhanced Decohesion

The HEDE model, advanced principally by Oriani and Josephic, proposes that hydrogen segregated at the crack tip or along grain boundaries reduces the theoretical cohesive strength σ_c of the metallic bond. The cohesive energy γ₀ in the absence of hydrogen is reduced to an effective value γ_H that depends on the local hydrogen chemical potential μ_H:

γẋ = γ₀ + (∂γ / ∂μẋ) × Δμẋ

When Δμẋ > 0 (enrichment above equilibrium), γẋ < γ₀
→ Critical stress intensity Kᴰ decreases without change in plastic zone size

HEDE predicts predominantly intergranular (IG) or cleavage-like transgranular fracture morphology. TEM studies of hydrogen-charged ferritic steel have confirmed hydrogen segregation at prior austenite grain boundaries (PAGBs) and lath martensite boundaries, consistent with HEDE. The mechanism operates at the crack tip ahead of the plastic zone and does not require dislocation motion — this explains why very high-strength steels that are already heavily dislocation-depleted can fracture with minimal macroscopic ductility.

HELP — Hydrogen-Enhanced Localised Plasticity

The HELP mechanism, developed by Birnbaum and Sofronis, proposes the opposite of HEDE at the dislocation scale. In situ TEM observations under hydrogen gas confirmed that hydrogen accelerates dislocation motion by reducing the elastic interaction energy between dislocations and between dislocations and obstacles (precipitates, solutes). The shear stress to move a dislocation past an obstacle is reduced by a factor proportional to the hydrogen concentration:

Δτ = τ₀ × [1 − cẋ(r) × ΔV / (kᴰT)]

where:
  τ₀  = obstacle strength without hydrogen
  cẋ(r) = local hydrogen concentration at dislocation
  ΔV  = activation volume for obstacle bypass
  kᴰT  = thermal energy

The increased dislocation mobility causes intense slip localisation in narrow bands immediately ahead of the crack tip, while remote regions remain nearly undeformed. This strain localisation eventually produces a ductile micro-void mechanism on the microscale, but the macroscopic result is a dramatic reduction in fracture toughness, notch ductility, and fatigue crack growth threshold. Fractographs of HELP-dominated fractures often show shallow dimples and transgranular features — superficially “ductile” in appearance but occurring at stress intensities far below K_Ic in the hydrogen-free case.

AIDE — Adsorption-Induced Dislocation Emission

A third mechanism, AIDE (Lynch, 1988), proposes that hydrogen adsorbed at the crack tip surface weakens interatomic bonds sufficiently to facilitate dislocation nucleation and emission from the crack tip itself. AIDE can occur at stress intensities below the threshold for stable crack growth by classical fracture mechanics. It produces characteristic fractographic features including small dimples and evidence of plasticity confined to within a few nanometres of the fracture surface. AIDE is considered most relevant for electrochemical conditions where surface coverage of adsorbed hydrogen is high.

Hydrogen pressure — blistering and HIC

Where two hydrogen atoms meet at a vacancy, void, or inclusion interface, they can recombine to form H₂ gas. Because molecular hydrogen cannot diffuse back through the steel lattice, internal H₂ pressure builds. At planar inclusions (particularly elongated MnS) in hydrogen-rich pipe steels, this pressure can reach several hundred MPa — exceeding the local yield strength and fracturing the steel in a plane parallel to the rolling direction. The resulting laminar cracks grow and link stepwise between adjacent planes: this is Hydrogen-Induced Cracking (HIC), also called stepwise cracking or blister cracking. The HIC susceptibility of pipeline steel is evaluated per NACE TM0284, with acceptance criteria based on Crack Length Ratio (CLR), Crack Thickness Ratio (CTR), and Crack Sensitivity Ratio (CSR).

Hydrogen-Assisted Cracking (HAC) in Welded Structures

In welded structural and pressure-vessel steels, HAC (also called cold cracking or delayed cracking) occurs when three conditions are met simultaneously — the Granjon triangle:

• Susceptible microstructure: martensite or upper bainite with hardness above approximately 250–300 HV is most at risk. The higher the carbon equivalent (CE), the greater the hardenability and susceptibility.
• Sufficient hydrogen: diffusible hydrogen above the threshold for the steel’s strength level (as low as 1–2 mL/100g for high-strength grades).
• Tensile stress: residual welding stress, applied stress, or both. Triaxial stress states (thick sections, restrained joints) are most dangerous because they raise the hydrostatic stress component and therefore the hydrogen fugacity at microstructural damage sites.

HAC typically initiates within 48–72 hours after welding as the weld cools and residual stresses develop. This incubation period reflects the time for hydrogen to diffuse from the weld metal to the HAZ and accumulate at susceptible microstructural features. Cracking almost always occurs in or near the coarse-grained HAZ (CGHAZ) where martensite is hardest and grain size largest. Hydrogen-induced cracking in welds is mitigated by preheat, inter-pass temperature control, low-hydrogen consumables, and post-weld heat treatment (PWHT).

Carbon equivalent and preheat estimation

The International Institute of Welding (IIW) carbon equivalent CE_IIW and the Ito–Bessyo parameter P_cm are used to predict HAC susceptibility and required preheat temperature T_p:

CEᴮᴮᴼ = C + Mn/6 + (Cr+Mo+V)/5 + (Ni+Cu)/15

Pcm = C + Si/30 + (Mn+Cu+Cr)/20 + Ni/60 + Mo/15 + V/10 + 5B

Graville preheat (simplified Lorenz–Degenkolbe):
Tᵖ (°C) = 697 × C✛ − 254 (for CE > 0.45, HD = 5 mL/100g)

where C✛ is the carbon equivalent from the applicable formula.

High-strength structural steels commonly require preheat temperatures of 100–200°C for multi-pass groove welds; pipeline girth welds in Grade X70 may require no preheat in favourable conditions but demand low-hydrogen consumables (HD ≤ 5 mL/100g per AWS A5.01 Level H5) and dry electrode storage.

Refer to our detailed guide on HAZ microstructure evolution for further treatment of CGHAZ and ICHAZ transformation behaviour.

High-Strength Steels and Special Susceptibility Cases

Ultra-high-strength steels (UHSS) and fasteners

Grade 12.9 fasteners and similar components with tensile strengths above 1200 MPa are extremely susceptible to hydrogen embrittlement from electroplating, acidic cleaning, and in-service sour environments. The threshold stress intensity K_Ith for hydrogen-assisted cracking in 1400 MPa martensite can be below 20 MPa√m — compared to K_Ic > 60 MPa√m without hydrogen. ASTM F519 and ASTM F3125 govern HE testing of fasteners. ISO 9588 specifies mandatory post-plating bake-out procedures.

Duplex stainless steel and HISC

Duplex stainless steels (DSS, e.g., SAF 2205 / UNS S31803) and super-duplex grades (SDSS, e.g., SAF 2507 / UNS S32750) have excellent corrosion resistance but are susceptible to Hydrogen-Induced Stress Cracking (HISC) when subjected simultaneously to: cathodic potentials below −0.80 V_Ag/AgCl, applied or residual tensile stress above approximately 80% of the 0.2% proof strength, and significant component section thickness (>40 mm). Hydrogen enters the ferrite phase preferentially (D_H,ferrite ≈ 10⁻¹²–10⁻¹¹ m²/s in DSS at ambient, much slower than in plain ferritic steel due to austenite barriers) and causes intergranular cracking along ferrite–austenite boundaries. DNVGL-RP-F112 provides material qualification and design guidelines including derating of allowable stresses as a function of section thickness and CP potential.

Pipeline steels and sour service (HIC/SSC)

In oil and gas pipelines carrying wet H₂S, two distinct failure modes apply:

• HIC (Hydrogen-Induced Cracking): pressure-driven laminar cracking at inclusions, evaluated per NACE TM0284 in H₂S-saturated 5% NaCl + 0.5% acetic acid (Solution A). Mitigation requires calcium treatment to globularise sulphide inclusions, low sulphur (≤0.002% S in EW grade), low carbon equivalents, and homogeneous microstructure without banding.
• SSC (Sulphide Stress Cracking): a form of hydrogen embrittlement that occurs in hard heat-affected zones and high-strength parent metal under applied or residual tensile stress in wet H₂S environments. ISO 15156-2 limits maximum hardness to 250 HV Vickers (22 HRC) for carbon and low-alloy steels in sour service; weld HAZ hardness must meet this limit under production WPS qualification.

See our article on corrosion mechanisms and pitting corrosion for complementary electrochemical background.

Detection, Testing, and Measurement

Electrochemical permeation (ASTM G148 / ISO 17081)

The Devanathan–Stachurski dual-cell technique measures hydrogen permeation through a steel membrane. One cell charges hydrogen electrochemically; the other detects the oxidation current as hydrogen emerges at the opposite surface. The method yields D_eff, the sub-surface concentration C₀, and the trap site density N_T from the steady-state permeation flux J_∞ and the breakthrough time t_lag:

D✛?? = L² / (6 × tᴸₐₐ)
J∞ = Dⁿ × C₀ / L
Nₜ (approx.) = (Dⁿ⁄D✛?? − 1) × Nⁿ

where L = membrane thickness (m)

Diffusible hydrogen in weld metal (ISO 3690)

The standard method for quantifying diffusible hydrogen in deposited weld metal: a weld bead is quenched immediately after welding, sealed, and incubated at 45°C for 72 hours with evolved gas collected by mercury displacement or gas chromatography. Results in mL H₂/100g weld metal are used to classify consumables: HD ≤ 5 mL/100g = H5 (low hydrogen), HD ≤ 3 mL/100g = H3 (very low hydrogen) per AWS A5.01 / EN ISO 2560. Most major fabrication codes (ASME IX, EN 1011-2, AS/NZS 2980) mandate H5 or H3 consumables for high-strength steel welds.

Our guide to Charpy impact testing and hardness testing covers complementary mechanical qualification methods routinely used alongside hydrogen testing in WPS qualification.

Slow strain rate testing (SSRT, ASTM F1624 / ISO 7539-7)

A pre-charged or in-situ charged specimen is pulled to fracture at a very low strain rate (typically 10⁻⁶–10⁻⁵ s⁻¹) in a controlled environment. Embrittlement is quantified by the Index of Hydrogen Embrittlement (IHE) or reduction in ductility:

IHE = 1 − (RAẋ / RA₀)

where:
  RAẋ = reduction in area in hydrogen environment
  RA₀ = reduction in area in inert environment (Ar or vacuum)

Values >0.1 (10% degradation) indicate significant susceptibility.

K_Ith determination — threshold stress intensity for hydrogen cracking

Pre-cracked compact tension (CT) or wedge-opening load (WOL) specimens are statically loaded or incrementally loaded in a hydrogen-containing environment to find the minimum stress intensity below which crack growth does not occur. ASTM F1624 (step-loading method) and ISO 7539-6 (constant load, pre-cracked specimen) are the primary protocols. K_Ith values for high-strength steels in hydrogen service range from <15 MPa√m (tempered martensite, 1400 MPa) to >40 MPa√m (high-toughness, low-sulphur grades at ≤900 MPa).

Mitigation Strategies

Material selection and microstructural design

• Limit maximum hardness and strength in sour service (≤250 HV per ISO 15156-2).
• Specify low sulphur (≤0.002%), calcium treatment, and low carbon equivalent for HIC-resistant pipeline steel per NACE TM0284 evaluation.
• Select tempered lower bainite rather than tempered martensite at equivalent strength: lower bainite offers better resistance to hydrogen embrittlement due to finer effective grain size and more uniformly distributed carbides.
• Use austenitic stainless steel, nickel alloys (Alloy 625, 825), or titanium alloys where hydrogen service is unavoidable; their FCC/HCP structures dramatically reduce hydrogen diffusivity.
• Incorporate fine TiC or NbC precipitates in alloy design to create high-density reversible/irreversible trap fields that retard hydrogen diffusion to crack-susceptible sites.

Process control — welding

• Use low-hydrogen consumables (HD ≤ 5 mL/100g or lower per code) and store electrodes in heated ovens (100–150°C) per manufacturer recommendation.
• Apply preheat and maintain minimum inter-pass temperature per CE_IIW or P_cm calculation or as specified in the WPS.
• Post-weld hydrogen bake-out at 250–300°C for 1–2 hours before cooling, where PWHT is not immediately practicable: this accelerates hydrogen diffusion out of the HAZ and reduces the risk of delayed cracking during the 48–72 hour critical window.

Post-process bake-out for electroplated components

For electroplated high-strength steel fasteners and components (grade 10.9, 12.9, springs, aircraft hardware), bake-out removes absorbed hydrogen before embrittlement can cause delayed fracture under assembly preload. ASTM B633 and ISO 9588 specify:

Bake-out temperature: 190–220°C
Holding time:         minimum 8 hours
Process window:       begin within 4 hours of plating (before oxide sealing)
Re-plating:           bake-out must be repeated after any re-electroplating

Avoid >250°C for tempered martensite (Grade 12.9) to prevent over-tempering
and hardness loss (<2 HRC reduction limit per ISO 898-1).

Surface coatings and barriers

Electroless nickel coatings, physical vapour deposition (PVD) TiN/CrN, and aluminium-based conversion coatings can reduce hydrogen entry flux by increasing surface recombination or reducing cathodic current density. None completely prevents hydrogen ingress under severe electrochemical conditions. For subsea cathodic protection, sacrificial zinc or aluminium anodes should be designed to maintain potential above the HISC risk threshold (−0.80 V vs Ag/AgCl for SDSS) without overprotection. Refer to our coverage of corrosion protection coatings for further detail.

Standards Summary

Industrial Applications and Case Studies

Offshore subsea production equipment

Super-duplex stainless steel manifolds and forged connectors operating on cathodic protection are routinely qualified against HISC per DNVGL-RP-F112. Maximum allowable stresses are derated to 80–100% of the 0.2% proof strength for sections below 25 mm, decreasing to 50–60% for sections above 100 mm. Qualification testing per ISO 15156-3 Annex B includes U-bend and C-ring specimens charged at −1.05 V_Ag/AgCl for 30 days at 23°C. See also our article on grain boundary structure and segregation for how boundary character affects hydrogen embrittlement resistance.

High-pressure hydrogen storage vessels

Type III and Type IV composite overwrapped pressure vessels (COPV) for hydrogen fuel cell vehicles (H₂ at 70 MPa) use aluminium or plastic liners precisely to avoid steel–hydrogen compatibility issues. Where steel is used in valve bodies and fittings, austenitic Grade 316L or nickel-based Alloy 718 is specified. The SAE J2579 and AIAG standards govern materials qualification, requiring testing at −40°C and elevated H₂ pressures to confirm that K_Ith in gaseous H₂ environment exceeds the maximum stress intensity generated under burst conditions with a specified safety factor.

Automotive press-hardened steel (PHS)

Ultra-high-strength 22MnB5 boron steel (Usibor, ZM grades) hot-stamped to 1500–2000 MPa is used in A/B-pillars and door intrusion beams. Delayed fracture from hydrogen absorbed during cataphoresis coating (e-coat) is a primary development challenge for ≥1800 MPa grades. Testing to VDA 238-300 (hole expansion with delayed fracture assessment) and accelerated H-charging followed by SCC testing quantify hydrogen sensitivity. Alloy additions of Ti and Nb to create fine carbide traps and low-temperature tempering to 200°C post-coiling are used to manage susceptibility. Martensite formation and morphology and quench and tempering responses are directly relevant to optimising PHS resistance.

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