Ultra High Tensile Strength Stainless Steel (2023)

January 11, 2023

Keywords: ultra-high strength stainless steel, strengthening and toughening mechanism, hydrogen embrittlement, stress corrosion, precipitated phase, reverse transformed austenite

Application of High Strength Stainless Steel

High-strength stainless steel is widely used in aerospace, marine engineering and energy fields, such as:

  • The main bearing member of the aircraft
  • fastener
  • satellite gyroscope
  • spaceship shell
  • Offshore oil platform
  • auto industry
  • nuclear energy industry
  • Gear and Bearing Manufacturing

The development history of high-strength stainless steel

  • In order to meet the needs of aerospace and marine engineering for high-performance corrosion-resistant structural steel, the American Carnegie Illionois Steel Company successfully developed the first generation of martensitic precipitation-hardening stainless steel – Stainless W in 1946.
  • On the basis of the Stainless W steel alloy system, Cu and Nb elements are added and Al and Ti elements are removed. Arm-co Steel Company of the United States developed 17-4PH steel in 1948. Because of its good strength, toughness and corrosion resistance, it is not only used in F-15 aircraft landing gear components, but also widely used in the manufacture of fasteners and engines. parts, but its cold deformation ability is poor. In order to reduce the high-temperature δ-ferrite that is unfavorable to the transverse mechanical properties, by reducing the content of ferrite-forming element Cr and increasing the content of Ni element, 15-5PH steel was developed, which overcomes the transverse ductility of 17-4PH steel Poor shortcomings, has been used in the manufacture of ships and civil aircraft and other load-bearing components.
  • In the early 1960s, the International Nickel Corporation invented maraging steel and introduced the concept of maraging strengthening for the development of high-strength stainless steel, thus opening the curtain of the development of maraging stainless steel.
  • In 1961, American Carpenter Technology Company first developed the Mo-containing maraging stainless steel Custom450.
  • In 1967 and 1973, Pyromet X-15 and Pyromet X-12 were developed successively. During this period, the United States also successively developed AM363, In736, PH13-8Mo, Unimar CR, etc.
  • Martin et al obtained the invention patents of Custom465 and Custom475 steels in 1997 and 2003 respectively, and applied them in civil aviation aircraft.
  • The UK has developed high-strength stainless steel grades such as FV448, 520, 520(B), and 520(S).
  • Germany developed Ultrafort401, 402 and so on in 1967 and 1971.
  • In addition to imitating and improving American steel grades, the former Soviet Union also independently researched a series of new steel grades. Common steel grades include 0Х15Н8Ю, 0Х17Н5М3, 1Х15Н4АМ3, 07Х16Н6, etc., as well as steel grades with higher Co content, such as 00Х12К14Н5М5Т, 00Х14К14Н4М3Т, etc.
  • In 2002, QuesTek of the United States undertook the pollution prevention project of the Strategic Environmental Research and Development Program (SERDP) of the US Department of Defense. Through the Materials Genome Project, it designed and developed a new type of ultra-high-strength stainless steel Ferrium® S53 for aircraft landing gear, and published it at the end of 2008. AMS5922 Aerospace Standard, Ferrium®S53 has a strength of about 1930 MPa and a fracture toughness (KIC) of more than 55 MPa m1/2. It was added to the MMPDS backbone material manual in the United States in 2017. This material has been successfully applied to A-10 in the United States. Fighter planes and T-38 aircraft are the preferred materials for the landing gear of the next generation of carrier-based aircraft.

Research progress of ultra-high strength stainless steel

The good properties of ultra-high-strength stainless steel mainly include ultra-high strength, excellent plasticity and toughness, excellent corrosion resistance, stress corrosion resistance and corrosion fatigue performance.

The following is the progress of exploring these properties of ultra-high-strength stainless steel.

Alloy Design and Strengthening Phases in High-Strength Stainless Steel

Typical room temperature structures of ultra-high-strength stainless steel include:

1. Fine lath martensite matrix

Lath martensite has high strength due to its own high dislocation density.

2. Appropriate amount of residual (or reverse transformation) austenite

Metastable residual (reverse transformation) austenite can relieve the stress concentration at the crack tip and improve the toughness of the material.

3. Precipitation strengthening phase dispersedly distributed

The nano-scale strengthening phase precipitated during the aging treatment can further improve the strength of the steel. According to the alloy composition of the precipitated phase, it can be divided into three categories, namely carbide (MC, M2C), intermetallic compound (NiAl, Ni3Ti) and element Enriched phase (ε phase, α’ phase), etc. In ultra-high-strength stainless steel, the strengthening potential of the precipitated phase depends on the nature of the precipitated phase and its size, number density, volume fraction, and spatial distribution. Whether the optimal performance can be obtained mainly depends on the control of the thermal and kinetic characteristics of the precipitation behavior of the precipitated phase, and then guides the regulation of the alloy composition and the formulation of the heat treatment process.

Research on the relationship between chemical composition and mechanical properties


When designing the composition of ultra-high-strength stainless steel, in order to ensure that the steel has good corrosion resistance, the content of Cr in general steel should be greater than 10%, and Cr is also an element that reduces the martensitic transformation temperature.


Ni can improve the potential and passivation tendency of stainless steel, increase the corrosion resistance of steel, improve the plasticity and toughness of steel, especially the toughness of steel at low temperature, and Ni will also form a strengthening η-Ni3Ti phase.


The addition of Mo is mainly to increase the secondary hardening effect. About 2% Mo can make the steel maintain a high hardness under different solution treatment conditions, and the Mo-rich precipitates precipitated during the aging process play a strengthening role. Make steel maintain good toughness, and Mo can also improve the seawater corrosion resistance of stainless steel.


Co can inhibit the recovery of dislocation substructure in martensite, provide more nucleation sites for the formation of precipitates, reduce the solubility of Mo in α-Fe, and promote the formation of Mo-containing precipitates.


Adding a small amount of Ti to the steel will significantly increase the strength of the steel, but excessive addition will reduce the toughness of the steel.

The chemical composition and mechanical properties of typical ultra-high-strength stainless steel are shown in the following chart:

15-5PH steel

As a typical representative of the first generation of high-strength stainless steel, the alloying characteristics of 15-5PH steel are:

  • About 15% Cr is used to ensure the corrosion resistance of steel;
  • The Ni content of about 5% can balance the Cr-Ni equivalent of the steel used in the experiment, so that the steel can obtain a martensitic structure at room temperature, and at the same time reduce the δ-ferrite in the steel;
  • Adding about 4% Cu plays a strengthening role;
  • A small amount of Nb can form MC phase with C, which plays the role of pinning grain boundaries and refining grains.
  • After aging treatment at 550 ℃, a large number of Cu-rich phases with fcc structure precipitated on the martensite matrix, and the orientation relationship between the Cu-rich phase and the martensite matrix satisfies the K-S relationship (111)Cu//(011)M, [11ˉ0] Cu//[11ˉ1]M.

Studies by Habibi-Bajguirani et al. have shown that there are two different types of Cu precipitates in 15-5PH steel during the aging process. When aging below 500 ℃, cluster particles with bcc structure will be formed first. This cluster It will subsequently evolve into a 9R structure, and finally transform into an fcc precipitated phase. The X-ray microanalysis results of the precipitated phase extract show that this precipitated phase is actually a Cu-rich phase. When aging at 650~700 ℃, the fcc Cu-rich phase maintains a coherent relationship with the matrix at first, and then transforms into a semi-coherent K-S relationship.


As a typical representative of the second-generation high-strength stainless steel, PH13-8Mo adopts a low-carbon alloy design, and its characteristics are:

  • About 13% Cr is used to ensure the corrosion resistance of steel;
  • About 8% Ni can make up for the Cr-Ni equivalent imbalance in the Schaeffler diagram caused by low carbon, reduce the δ-ferrite content, and make the steel obtain lath martensitic structure;
  • Adding 1% Al can form a strengthening phase in the steel and play a role in strengthening the matrix.

Schober et al. studied the effect of Ti element on the evolution of precipitates during the aging process:

  • In the PH13-8Mo steel without adding Ti element, the precipitated phase is only NiAl phase.
  • After adding Ti element, the precipitated phases in the steel are G phase and η phase. The orderly intermetallic compound NiAl is precipitated in the PH13-8Mo steel without adding Ti element at the initial stage of aging treatment. With the prolongation of aging time, the alloying elements in the NiAl phase gradually tend to the stoichiometric equilibrium and the hardness reaches the maximum value. In the steel with Ti added, a precipitation phase rich in Ni, Si, Al and Ti is precipitated in the steel at the initial stage of aging treatment, and the hardness of the steel reaches the maximum at this time. With the prolongation of aging time, ellipsoidal Ni16Si7Ti6-G phase and short rod-shaped Ni3(Ti, Al)-η phase will be formed in the steel.


Li et al. studied a Cr-Ni-Co-Mo-based martensitic precipitation-hardening stainless steel with a strength of up to 1900 MPa, and believed that the ultra-high strength was obtained due to the composite strengthening of multiple strengthening phases.

The nominal composition of the steel is 0.004C-13.5Cr-12.7Co-3.3Mo-4.4Ni-0.5Ti-0.2Al (atomic fraction %).

There are mainly three kinds of precipitated phases in steel, η-Ni3(Ti, Al) phase, Mo-rich R’ phase and Cr-rich α’ phase. These precipitated phases are transformed from Ni-Ti-Al-rich, Mo-rich and Cr-rich cluster particles at the early stage of aging respectively. During the aging process, the η-Ni3(Ti, Al) phase grows slowly due to the segregation of the Mo-rich R’ phase and the Cr-rich α’ phase.

A new calculation model for alloy design

From the perspective of the development of high-strength stainless steel, as the strength level increases, the strengthening of a single strengthening phase gradually develops into multi-phase composite strengthening. Compared with the strengthening of a single type of precipitated phase, composite strengthening is more conducive to the further improvement of steel strength.

However, the influence of alloy composition and aging system on the precipitation and growth behavior of different types of precipitated phases is quite different. Considering that different alloy compositions and heat treatment systems can obtain different and various precipitated phases when designing new steel grades, there are still deficiencies in the alloy design process using traditional trial-and-error experiments and artificial neural network simulations based on data accumulation. A new type of physical metallurgy-based model is urgently needed.

For example, Xu et al. and Parn et al. proposed a machine learning-based calculation model for alloy composition. This model integrates the alloy composition and corresponding heat treatment parameters, enabling the desired properties to evolve within a genetic framework. This model is applied to the design of ultra-high-strength steel with MC carbide as the strengthening phase. It is also suitable for Cu clusters, Ni3Ti, and NiAl precipitated phases. It can also be applied to design a multi-type strengthening phase, including MC carbide, rich Cu phase and Ni3Ti intermetallic compound strengthen the alloy together. The model includes the simulation of corresponding parameters such as steel mechanical properties, corrosion resistance and microstructure, which provides a more reliable path for alloy composition design.

Toughening Phase and Toughening Mechanism

The effect of reverse transformed austenite on the toughness of high-strength stainless steel is closely related to its morphology, content, dispersion and stability.

Its characteristics are affected by the heating rate, isothermal temperature and time of the heat treatment process, the diffusion and segregation of austenite forming elements, the nucleation position and size of austenite, and the dislocation density in the matrix.

Existing studies have shown that there are three mechanisms for the formation of reverse transformed austenite,

  • Diffusion-free shear inversion mechanism,
  • variant restriction mechanism,
  • Retained austenite growth mechanism.

The shearing mechanism originates from the inverse process of the non-diffusion shearing mechanism from austenite to martensite. The reverse transformed austenite formed by martensite that maintains a certain crystal degree phase relationship with the original austenite, and the original austenite maintain the same phase relationship.

The modification restriction mechanism points out that during the formation of reverse transformed austenite controlled by diffusion, its nucleation position will strictly maintain a certain crystallographic phase relationship with the original austenite, carbide and matrix, thus limiting the transformation of reverse transformed austenite. Types of variants. The growth mechanism of retained austenite believes that the residual austenite in martensitic steel after quenching will continue to grow through the diffusion of austenite stabilizing elements in the subsequent tempering process, thereby further “reversing transformation”. “For the new austenitic organization.

Research on 0Cr13Ni4Mo martensitic stainless steel shows that carbide (Cr23C6) and austenite co-precipitate during tempering in the two-phase region slightly higher than the austenite transformation start temperature (AS). Further analysis of carbide, austenite and the distribution of Cr and Ni elements on the interface shows that the segregation of Cr in carbide promotes the distribution of Ni element to reverse austenite, and the enrichment of Ni element reduces reverses the chemical driving force for austenite formation and increases the interfacial energy,

Therefore, the Ni-rich region can be used as the nucleation site of reverse transformed austenite during tempering, that is, the formation of reverse transformed austenite is controlled by the diffusion of Ni element.

Further increase the tempering temperature, although the diffusion of atoms is more significant, but due to the increase in temperature, the driving force condition for the transformation of tempered martensite to austenite has been satisfied, so the formation mechanism of reverse transformed austenite at this time is no The shear mechanism of diffusion.

In order to further explain the modification restriction mechanism, Nakada et al. studied the crystal degree phase relationship between reverse transformed austenite and prior austenite and martensite matrix. After tempering of 13Cr-6Ni steel, in an original austenite grain, the reverse transformed austenite is not only uniformly distributed in the martensite lath boundary, but also has reverse transformation at the interface between blocks and packets. Transformed austenite, and most of them maintain the same orientation as the original austenite, while a small part of the orientation is different from the original austenite. There may be 12 phase relationships of reverse transformed austenite variants in a prior austenite habit surface and a martensite lath group.

It can be observed that under the premise of following the K-S relationship, there are only 6 different directions of martensite lath bundles parallel to the close-packed plane, and there are only 2 reverse transformation austenite bundles inside each martensite lath bundle. body variant.

This shows that due to the triple symmetry of austenite in the {111} γ plane family, the 12 reverse-transformed austenite variants in a martensitic lath group can be divided into 2 types, that is, the same as the original austenite Oriented V1 variants and V2 variants that are twinned to V1.

According to the two-dimensional construction model proposed by Lee and Aaron-son, the critical nucleus shape of reverse transformed austenite should meet the requirement of minimizing the nucleation energy.

The reverse transformed austenite formed at the lath interface is usually consistent with the orientation of the original austenite grains, and the α’/γ interface of the core maintains the K-S relationship with both sides of the martensite matrix, while the original austenite grain boundary The austenitic core only maintains the K-S relationship with the matrix on one side.

Therefore, the reverse transformed austenite at the original austenite grain boundary will form a spherical shape due to being wrapped by coherent and incoherent interfaces, and the difference in surface energy and elastic strain energy at the two sides of the boundary, while at the lath The reverse-transformed austenite tends to form elongated needle-like morphology.

The increase of the reverse transformed austenite content can improve the plasticity and toughness of the material, while too much reverse transformed austenite often leads to the deterioration of the yield strength of the steel.

Schnitzer et al. respectively calculated the influence of strengthening phase NiAl and toughening phase reverse transformation austenite on the overall yield strength in PH13-8Mo, and the 40% decrease in yield strength after aging treatment was attributed to the high content of reverse transformation austenite, The rest is attributed to the coarsening of the NiAl phase.

Therefore, in the case where high toughness is required, a higher aging temperature should be used to increase the reverse transformed austenite content, but at the cost of losing the strength of the material. In addition, some studies have also found the adverse effect of reverse transformed austenite on plasticity. For example, the results of Viswanathan et al. showed that the improvement of plasticity by reverse transformed austenite only occurs in the early stage of aging, and the prolonged time will also cause serious brittleness of the material. fracture.

Susceptibility to Hydrogen Embrittlement and Stress Corrosion Research

As the strength level increases, high-strength steels become more sensitive to stress corrosion cracking (stress corrosion cracking, SCC) and hydrogen embrittlement (hydrogen embrittlement, HE). In particular, when polluting or corrosive gas components and H atoms act on high-strength steel in combination with stress, it is very easy to cause crack initiation and gradually expand until cracking.

This kind of fracture is the main failure mode of high-strength steel structural parts serving in corrosive environments, causing huge safety hazards and property losses.

Susceptibility to hydrogen embrittlement

Diffusible hydrogen is the main factor causing the plasticity loss of steel. Any measure that reduces the mobility of diffusible hydrogen can effectively improve the resistance to hydrogen embrittlement susceptibility of materials.

Strong hydrogen traps can significantly increase the content of supersaturated hydrogen absorbed by steel, thus making the hydrogen entering the matrix harmless.

The above point of view has been confirmed to a certain extent in the observation of hydrogen-induced delayed fracture of high-strength steel, that is, when high-strength steel is under the action of a static stress lower than its tensile strength, it will undergo instantaneous brittle fracture after a period of service. The failure under static load is due to the intrusion of H atoms into the matrix.

As the main strengthening phase and toughening phase in steel, a large number of dispersed second phase strengthening particles and reverse transformed austenite precipitated during aging can be regarded as important hydrogen traps in steel.

A lot of research has focused on regulating the number and density of “benign hydrogen traps” (benignhydrogen traps) in steel through heat treatment to prevent the diffusion of H in the material, thereby improving the material’s resistance to hydrogen embrittlement sensitivity.

A large number of studies have shown that carbides are typical “benign hydrogen traps” in steel and can effectively increase the hydrogen embrittlement susceptibility of steel. For example, by spheroidizing cementite particles or refining cementite by rapidly heating to tempering temperature after forming and cooling in the austenite single-phase region, the hydrogen embrittlement susceptibility resistance of steel can be effectively improved.

In addition, by adding microalloying elements such as Ti, V, and Nb, carbides such as TiC, VC, and NbC are formed in the steel, which can be used as effective hydrogen traps. Takahashi et al. used APT to directly observe that TiC and V4C3 traps captured deuterium atoms. H is mainly trapped on the interface between TiC and the matrix, while the trap sites in V4C3 are mainly the core positions of misfit dislocations on the semi-coherent interface. With the help of first-principles calculations and finite element analysis, it is further confirmed that for TiC precipitation, the TiC-matrix interface is the main hydrogen trap, while the carbon vacancies are the main trap sites in V4C3.

There are few reports on intermetallic compounds and element-rich phases as hydrogen traps.

Recently, Li et al. compared the hydrogen embrittlement behavior of 17-4PH steel and PH13-8Mo steel for the last stage of steam turbine blades. The research results showed that the type of precipitates in the steel and the crystallographic relationship between the martensite matrix and the precipitates, It is the main reason that PH13-8Mo steel has higher apparent hydrogen diffusion coefficient and lower apparent hydrogen solubility than 17-4PH steel.

Compared with the coherent β-NiAl phase in PH13-8Mo steel, there is a Cu-rich phase incoherent with the matrix in 17-4PH steel, which has a stronger ability to capture H atoms. This is because the radius of the octahedral gap of the Cu-rich phase is 0.0529 nm, which is about twice the radius (0.0206 nm) of the octahedral gap of the βNiAl phase.

Moreover, compared with the coherent interface between the β-NiAl phase and the matrix, the non-coherent interface between the Cu-rich phase and the matrix can trap more H atoms.

In addition, the core of the misfit dislocation on the coherent interface and the less distorted lattice adjacent to the core are weak hydrogen traps, and the hydrogen de-trapping energy of the incoherent precipitated phase is higher than that of the coherent dislocation. The desorption energy of the lattice precipitated phase.

Compared with the martensite matrix, the diffusion rate of H in the residual (or reverse transformation) austenite is lower (diffusion rate in austenite: 10-15~10-16m2/s, in martensite Diffusion rate: 10-10~10-12m2/s), and the solubility of H in austenite is higher than that in martensite. In addition, the pinning energy of austenite for H can reach 55kJ/mol, making it an irreversible H trap site.

However, the influence of austenite in steels of different systems relative to the hydrogen embrittlement susceptibility of the material is still widely debated. Some results show that the reverse transformed austenite and fine retained austenite in the steel can effectively prevent the diffusion of H in the matrix, thus improving the hydrogen embrittlement susceptibility resistance of the steel.

On the contrary, some scholars pointed out that the H atoms dissolved into the austenite can reduce its stacking fault energy, making the TRIP effect more likely to occur, and the new martensite as a “hydrogen source” will release H atoms, resulting in the material’s Brittle.

Fan et al. reported the effect of reverse transformed austenite on the hydrogen embrittlement fracture behavior of S41500 martensitic stainless steel (nominal composition 0.04C-13Cr-4.1Ni-0.6Mo-0.7Mn, %). In the reverse transformed austenite of Ni, there is no enrichment of H atoms at the interface of austenite/martensite and austenite/carbide.

The TEM observation results of the quasi-cleavage fracture of the sample after tempering treatment show that the fracture path is along the interface between the tempered martensite and the newly formed martensite (NFM) under the transformation-induced plasticity (TRIP) effect, which is Because most of the H has been captured by the reversed austenite instead of segregating at the original austenite grain boundary, which reduces the stability of the reversed austenite and promotes the martensitic transformation.

After the phase transformation occurs, the nascent martensite will act as a hydrogen source to release a large amount of H atoms, causing a large amount of H atoms to gather at the surrounding interface, and the resulting fracture morphology is a quasi-cleavage morphology rather than an intergranular fracture morphology.

Hydrogen-induced cracks generally nucleate at laths, isophase bundles, lath groups, and original Austrian grain boundaries, and then the cracks pass through the lath bundles under the action of external stress and propagate along the lath groups and original Austrian grain boundaries.

In high-strength stainless steel, many martensitic multi-level structure interfaces (original austenite grain boundary, martensite lath group boundary, martensite lath bundle boundary and martensite lath boundary) and phase boundaries are high-strength stainless steel. One of the reasons for the higher susceptibility to hydrogen embrittlement.

The research results of hydrogen diffusion and hydrogen embrittlement behavior in 17-4PH steel show that the hydrogen embrittlement susceptibility resistance of the solid solution state sample is higher than that of the peak aging state sample. This phenomenon is mainly due to the Cu-rich phase and matrix in the aging state sample. The phase boundary phase captures more H, and the weakening of the interfacial binding force causes the brittle fracture of the hydrogen-charged sample in the peak aging state.

With the increase of solution treatment temperature, the susceptibility to hydrogen embrittlement and hydrogen diffusion coefficient of 17-4PH steel first increased and then decreased.

This is mainly due to the effect of the solution temperature on the grain boundaries of the original austenite in the steel and the number density of the precipitated phases during the subsequent aging treatment. With the increase of the solution temperature, the original austenite grains become larger and the grain boundary area increases. decreases, but the solid solubility of the matrix for Cu atoms increases, which promotes the precipitation of Cu-rich phases during the aging process, and the increase in the density and size of the precipitated phases provides more phase interfaces, which together provide an interface that can trap H .

Obviously, the susceptibility to hydrogen embrittlement of high-strength stainless steel is jointly determined by the complex multi-level and multi-phase structure in the steel. Due to the limitations of analytical and characterization methods, it is still difficult to quantitatively determine the influence of various hydrogen traps on the hydrogen embrittlement susceptibility of high-strength stainless steel.

The influencing factors of hydrogen embrittlement susceptibility of high-strength stainless steels strengthened by different strengthening systems based on different strength levels still need to be studied systematically and deeply.

The susceptibility to hydrogen embrittlement of ultra-high-strength stainless steel with complex alloy system and multi-phase coupling strengthening needs to be studied urgently.

At present, the author’s team has developed a new type of 2200MPa high-strength stainless steel strengthened by multi-phase composite precipitation. ), the APT analysis results of the dual-aging sample are shown in the figure below.

It can be seen from the figure that there are obvious Mo/Cr/C, Mo/Cr and pure Cr-rich clusters in the steel. Further analysis shows that the precipitated phases in the steel include intermetallic compounds, carbides and Cr-rich phases. The strength is obtained by the coupled strengthening of three precipitates, and it is also the high-strength stainless steel with the highest strength level reported so far.

Stress corrosion cracking

The American aircraft component failure investigation report shows that stress corrosion cracking is one of the main forms of sudden failure accidents of key load-bearing components of aircraft during service.

Most of the landing gears are finally broken due to stress corrosion or fatigue crack growth.

At present, stress corrosion occurs not only in high-tech and industries such as aviation, aerospace, energy, and chemical industry, but also in almost all commonly used corrosion-resistant steels and alloys.

Therefore, analyzing the stress corrosion cracking mechanism of ultra-high-strength steel and the factors affecting stress corrosion of ultra-high-strength steel have great scientific value and practical significance for determining the stress-corrosion protection measures of ultra-high-strength steel.

The corrosion resistance of materials has become an important factor limiting stress corrosion cracking of high-strength steels, and pitting corrosion is the most common and most harmful form of corrosion.

Most stress corrosion cracking originates from pitting pits. During the aging treatment of ultra-high-strength stainless steel, precipitated phases precipitated from the supersaturated martensite matrix cause inhomogeneity in the microstructure. Primary source of pitting corrosion.

The passivation film near the precipitated phase is relatively weak, and the intrusion of Cl causes the destruction of the passive film, and a micro battery is formed between the precipitated phase and the matrix, thereby dissolving the matrix, exfoliating the precipitated phase, and forming pitting corrosion. For example, the Cr-rich carbide M23C6, M6C and the intermetallic compound Laves phase and σ are easy to form a Cr-poor area around them, resulting in the occurrence of pitting corrosion.

Luo et al. and Yu Qiang studied the effect of aging time on the microstructure and electrochemical behavior of 15-5PH ultra-high strength stainless steel using three-dimensional atom probe tomography.

Cu-rich clusters and (Cu,Nb) nanoparticles were observed when the aging time ranged from 1 to 240 min. Compared with the short-term aging treatment, the surface of the samples after long-term aging treatment was more susceptible to attack by Cl.

After aging for 240 minutes, the Cr content around the precipitates will also decrease, and these parts are prone to form Cr-poor areas. The reduction of the ratio of Cr/Fe in the passivation film is the reason for the decline of the pitting corrosion resistance of the passivation film.

In addition, the continuous precipitation of Cr-rich carbides on grain boundaries will reduce the intergranular corrosion resistance of steel. For example, studies have found that AISI316Ti stainless steel has higher resistance to intergranular corrosion than AISI321 stainless steel. The reason is that the precipitation of TiC reduces the formation of Cr-rich carbides, which are the precipitates that lead to intergranular corrosion. one of the things.

As the most important ductile phase in high-strength stainless steel, the content, morphology, size and stability of austenite will also affect the stress corrosion susceptibility of steel.

In the case of the same size, shape and stability, as the austenite content increases, the stress corrosion cracking threshold (KISCC) increases, and the stress corrosion cracking sensitivity of steel decreases.

The reason is that the film-like austenite structure formed on the martensitic lath boundary improves the toughness of the steel and reduces the growth rate of hydrogen-induced cracks. There are two main reasons for the decrease in the crack growth rate:

When the crack propagates from the martensite matrix to the film-like austenite, whether it continues to expand into the austenite or changes the direction of expansion to bypass the austenite structure, it will consume a lot of energy, resulting in a crack growth rate Reduced, increased stress corrosion resistance;

As mentioned above, H has higher solid solubility and lower segregation tendency in austenite structure, and the diffusion rate of H in austenite is much smaller than that in martensite structure, which is high The beneficial hydrogen traps in the high-strength stainless steel lead to a decrease in the susceptibility to hydrogen embrittlement at the front of the crack, which in turn reduces the crack growth rate and increases the stress corrosion susceptibility.

It should be noted that the stability of austenite is also a key parameter to determine the stress corrosion susceptibility of steel. After stress or strain induces martensitic transformation, the fresh martensite transformed from austenite cannot suppress the crack growth. It will also serve as a new source of hydrogen diffusion to increase the susceptibility to hydrogen embrittlement of steel.

In summary, the strength, toughness, stress corrosion and hydrogen embrittlement susceptibility of steel are all affected by the complex multi-level multi-phase structure, and the traditional trial and error method is used to design and manufacture ultra-high strength steel with both ultra-high strength, toughness and excellent service performance. Stainless steel is difficult, the cycle is long, and the cost is high.

Compared with the trial and error method, the rational design method, such as establishing a series of multi-scale analysis models of strength and toughness, stress corrosion performance and hydrogen embrittlement performance such as “atomic size-nanoscale-microscale”, will be more purposeful. Establish design standards for high-strength stainless steel through simulation analysis results, optimize the shape, size and content of precipitated phases, martensite and austenite structures in steel, and further combine multi-scale simulation with actual material development processes, which will greatly reduce Difficulty in material research and development, reducing cost input and shortening the research and development cycle.


As a metal structural material with excellent strength, toughness and service safety, high-strength stainless steel has broad application prospects in the fields of aviation, aerospace, ocean engineering and nuclear industry in the future.

In view of the harsh application environment of this type of steel, the exploration of a new generation of high-strength stainless steel should not only focus on further breaking through the bottleneck of ultra-high strength-excellent plasticity and toughness matching, but also take into account excellent service safety.

In the process of alloy design and heat treatment process formulation, the traditional trial and error method is gradually transitioned to rational design methods such as thermal/kinetic assisted alloy design, artificial intelligence machine learning, etc., so as to greatly improve the development cycle of new high-strength corrosion-resistant alloys, Save R&D costs.

The research on the strengthening and toughening mechanism of high-strength stainless steel still needs to be further in-depth, especially the understanding of the precipitation behavior of the second phase particles for multi-phase composite strengthening and the superposition of strengthening contribution values.

The research on the influence of austenite content, size, morphology and stability in steel on the toughness of high-strength stainless steel is relatively sufficient, but no effective mathematical model has been established to quantitatively estimate its contribution to the toughness of this steel.

In addition, the research on the stress corrosion cracking mechanism and hydrogen embrittlement susceptibility of ultra-high-strength high-strength stainless steel under complex strengthening system needs to be solved urgently, so as to provide a theoretical basis for the durability design of ultra-high-strength high-strength stainless steel.

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