Study on Cement Carbonation Resistance and Reinforcement in CCUS-EOR

Abstract

To investigate the mitigation of high-pressure CO₂-induced degradation of wellbore cement sheath in Carbon Capture, Utilization, and Storage–Enhanced Oil Recovery applications (CCUS-EOR), conventional Class G oil well cement and modified cement systems incorporating graphene, waterborne epoxy resin, and a composite of waterborne epoxy resin with graphene were formulated. This study presents the original comparative investigation on the long-term carbonation resistance of graphene-modified, waterborne-epoxy-modified, and their composite-modified oil well cements under 130 °C and 7 MPa CO₂ partial pressure, filling the research gap of unclear synergistic effects of the two modifiers in high-temperature CCUS environments. The specimens were subjected to simulated downhole conditions, and key properties, including compressive strength and permeability, were evaluated. The underlying mechanisms were elucidated through material characterization techniques such as X-ray diffraction, X-ray computed tomography, and scanning electron microscopy. Results indicated that the waterborne epoxy resin–modified cement system exhibited superior long-term carbonation resistance, achieving a 90 d compressive strength retention rate of 84%. The graphene-modified cement showed a 90 d compressive strength retention rate of 65%, while the waterborne epoxy–graphene composite system only retained 39.7% of its compressive strength at 90 d due to negative synergistic effects. The enhanced durability of the waterborne-epoxy-modified cement is attributed to the formation of a continuous polymeric film, which acts as a protective barrier against CO₂ penetration. This study provides valuable insights for the design of CO₂-resistant cement systems in CCUS-EOR environments.

1. Introduction

Driven by global carbon neutrality strategies, Carbon Capture, Utilization, and Storage–Enhanced Oil Recovery (CCUS-EOR) has become a vital technology for enhancing hydrocarbon production and reducing emissions [1]. However, prolonged exposure to high-concentration CO₂ during injection and production compromises wellbore integrity, as the cement sheath undergoes severe degradation, threatening long-term storage security [2,3]. Under high-pressure, high-temperature conditions, CO₂ penetrates cement pores in supercritical or aqueous form, reacting with hydration products such as calcium hydroxide (Ca(OH)₂) and calcium silicate hydrate (C–S–H) gel to form calcium carbonate (CaCO₃) and amorphous SiO₂·nH₂O. These reactions induce early-stage pore filling and densification but cause long-term pore structure deterioration, compressive strength reduction, permeability increase, and matrix embrittlement [4,5]. Furthermore, CO₂-induced corrosion weakens interfacial bonding at the cement-casing and cement-formation interfaces [6], promoting micro-annuli and leakage pathways [7]. Thus, developing advanced cement systems resistant to such harsh acidic environments is imperative.

To address CO₂ corrosion under high temperature and pressure, various strategies have been explored. Omosebi et al. [8] investigated how CO₂ concentration and pressure influence carbonation in Class G and H cements, noting that increased CO₂ levels generally raise compressive strength but reduce porosity and permeability, with pressure significantly accelerating degradation. Fakher et al. [9] developed a cement composite using fly ash and resin, demonstrating improved corrosion resistance under supercritical CO₂, though high fly ash content impaired slurry rheology and increased microcracking [10]. Zhang et al. [11] incorporated a soap-free emulsion with sodium styrene sulfonate and nano-SiO₂ into cement, reducing contact between hydration products and CO₂. The nano-SiO₂ consumed Ca(OH)₂, lowering its availability for carbonation. However, high-temperature CO₂ resistance remains underexplored, especially above 120 °C [12,13,14]. Studies by Yuan et al. [15], Peng et al. [16], and Bai et al. [17] on polymer-modified cement with corrosion-resistant additives (CRA) in H₂S-CO₂ environments revealed that film-forming polymers enhance acid resistance through sealing and pore-filling effects, informing design of CO₂-resistant slurries.

Graphene, known for its high aspect ratio and thermal conductivity, has been incorporated into polymers to improve mechanical properties, corrosion, and heat resistance [18]. Its sheet-like structure and nanoscale dimensions can fill cement voids and inhibit microcrack propagation. Epoxy resin is widely recognized for corrosion protection [19,20]. Waterborne epoxy (WER), which uses water as a dispersion medium, exhibits better handling safety and system compatibility than solvent-based epoxy [21,22], fully complying with green chemistry principles [23]. Compared to solvent-based epoxies, waterborne systems align better with field operations and green chemistry principles.

Existing studies separately verify the enhancement of graphene or waterborne epoxy on cement carbonation resistance, but few focus on their composite modification under high-temperature (≥130 °C) and high-pressure conditions. Graphene and waterborne epoxy are expected to produce synergistic reinforcement, yet their compatibility and synergistic mechanisms in harsh CCUS environments remain unclear. This study fills the gap by comparatively investigating graphene-modified, waterborne-epoxy-modified, and their composite-modified cement systems, revealing the failure mechanism of the composite material and the optimal modification path suitable for 130 °C high-temperature corrosion environments. Specifically, the carbonation resistance of these three cement systems was evaluated under aggressive conditions (130 °C, CO₂ partial pressure of 7 MPa, total pressure 25 MPa, salinity 22,863 mg/L). The evolution of cement properties and the underlying mechanism enhancing carbonation stability were systematically investigated, and the failure mechanism of the composite modification system was clearly elucidated. This study provides a reliable design basis for CO₂-resistant cement systems applied in high-temperature CCUS-EOR wellbore engineering.

2. Materials and Methods

2.1. Materials

Class G oil well cement, procured from Sichuan Jiahua Special Cement Co., Ltd., Leshan, China, is a widely used cementitious material in well cementing operations; its chemical composition is detailed in

Table 1. Chemical Composition of Class G Oil Well Cement (wt%).

CaO	SiO2	Fe2O3	Al2O3	MgO	SO3	Others
64.91	22.80	4.37	2.82	1.34	1.94	1.80

. Silica fume (industrial grade, SiO₂ ≥ 95 wt%), defoamer (polyether type), dispersant (sulfonated ketone-formaldehyde polymer), and retarder (organic phosphonate) were supplied by Dezhou Continental Shelf Petroleum Engineering Technology Co., Ltd., Dezhou, China). The silica fume serves to enhance the mechanical stability of the cement under high-temperature conditions and participates in a pozzolanic reaction with cement hydration products, thereby reducing the overall alkalinity of the system. The defoamer eliminates entrained air during slurry mixing. The retarder is employed to delay the hydration process, preventing premature setting. The dispersant, a ketone–formaldehyde polymer, improves the fluidity of the slurry. Graphene nanosheets were synthesized in-house by the liquid-phase exfoliation method [18], exhibiting a carbon content greater than 99.5%, a thickness of less than 3 nm, and a particle size ranging from 1.3~2.3 μm. Waterborne epoxy resin was prepared by the self-emulsification method, referred to in the literature as a white emulsion [21], with particle sizes between 0.2~0.5 μm and an epoxy equivalent weight of 500~2000 g/mol. Waterborne epoxy resin was prepared by the self-emulsification method [21], and it is a white emulsion with particle sizes between 0.2 and 0.5 μm and an epoxy equivalent weight of 500–2000 g/mol.

2.2. Preparation and Corrosion of Cement Samples

2.2.1. Preparation of Cement Slurry

Based on the formulation in

Table 2. Composition of Cement Slurry (mass ratio relative to cement, %).

Sample	Cement	Water	Retarder	Dispersant	Defoamer	Silica Fume	Graphene	Waterborne Epoxy Resin
PT	100	54	4	1	1	35	0	0
E8	100	46	4	1	1	35	0	8
G0.8	100	54	4	1	1	35	0.8	0
EG	100	46	4	1	1	35	0.8	8

, deionized water, retarder, dispersant, defoamer, and modifying materials were measured and combined into a homogeneous mixture. Graphene was ultrasonically dispersed for 10 min to prevent agglomeration before mixing with other components, then transferred into a mixing cup. Pre-weighed Class G oil well cement and 300-mesh silica fume were then introduced as a blended powder into the cup. All solid materials were added within 15 s. The cup was sealed and mixed at 12,000 rpm for 35 s using a constant-speed mixer (TG-3060 A, Shenyang Taige Petroleum Instrument Equipment Co., Ltd., Shenyang, China) to prepare the cement slurry.

The dosage of 8 wt% waterborne epoxy resin and 0.8 wt% graphene was determined based on preliminary optimization experiments and literature reports [18,21]. Within this dosage range, the cement slurry has good rheological properties and setting stability, avoiding performance deterioration caused by excessive resin agglomeration or graphene agglomeration. This study only investigates the single optimal dosage of each modifier, and the performance laws of other dosage ranges need to be further studied.

2.2.2. Preparation of Cement Samples

The prepared cement slurry was cast into API-standard square molds (50.8 × 50.8 × 50.8 mm³) and transferred to a high-pressure, high-temperature curing chamber (TG-7370D, Shenyang Taige Petroleum Instrument Equipment Co., Ltd., Shenyang, China). The chamber was filled with water and hermetically sealed. Curing proceeded at 130 °C and 25 MPa for 7 d. Owing to the size constraints of standard specimens, cylindrical samples (Ø25 mm × 50 mm) were drilled from the cured blocks for corrosion evaluation studies. The drilling process adopted low-speed water cooling to reduce micro-defects introduced by mechanical damage.

2.2.3. Corrosion of Cement Samples

The drilled cylindrical cement specimens were placed in a high-pressure, high-temperature autoclave (TL-3, Jingzhou Taling Machinery Co., Ltd., Jingzhou, China). Synthetic brine of specified salinity (composition in

Table 3. Composition of Simulated Saline Water (mg/L).

Ion Content
HCO3−	CO32−	Cl−	SO42−	Ca2+	Mg2+	Na+	K+
3112.02	60.02	11,166.75	119.11	372.74	67.82	3523.4	4441.5

) was added to fully submerge the specimens. The salinity of 22,863 mg/L is consistent with the formation water salinity of onshore CCUS-EOR reservoirs in Shengli Oilfield, which is used to reproduce the real downhole mineralization environment. Corrosion testing was conducted at 130 °C and 25 MPa total pressure (7 MPa CO₂ partial pressure). Temperature and pressure were continuously monitored and maintained throughout the experiment to ensure consistent test conditions.

2.3. Experimental Methods

2.3.1. Compressive Strength

The compressive strength of cement specimens was determined using a universal testing machine (HY-20080, Shanghai Hengyi Precision Instrument Co., Ltd., Shanghai, China) at a constant displacement rate of 1.2 kN/s. Triplicate specimens were tested for each condition, and the mean value was reported.

2.3.2. Gas Permeability

Specimens were oven-dried at 60 °C for 24 h until constant mass was achieved. The permeability and porosity of the dried cement samples were subsequently determined using a core flow apparatus (LDY 50-180, Nantong Yichuang Experimental Instrument Co., Ltd., Nantong, China). Triplicate measurements were performed, and the mean value was reported.

2.3.3. Corrosion Degree

The corrosion degree of cement is quantitatively defined as the ratio of the carbonated area to the total cross-sectional area. Cement hydrates create a highly alkaline environment (pH 12–13). Phenolphthalein solution was uniformly sprayed onto sectioned samples to distinguish the carbonated zone (colorless) from the uncarbonated zone (purple–red). Due to potential corrosion heterogeneity, the carbonated area ratio was quantified using digital image analysis (Equation (1), Figure 1) to evaluate the corrosion extent:

$$D = {\displaystyle \frac{S_1}{S_2}},$$

(1)

where S₁ represents the corroded cross-sectional area, S₂ denotes the original cross-sectional area, and D is defined as the corrosion degree, calculated as the ratio of S₁ to S₂.

2.3.4. Phase Composition Analysis

Corroded samples were cut into small pieces (≈5 mm × 5 mm × 5 mm), oven-dried at 60 °C for 12 h to remove free water, and ground into powder (<50 μm) for XRD and DTG testing. Phase composition changes in cement specimens before and after corrosion were characterized by X-ray diffraction (scanning range 5°~70° 2θ, rate 0.08°·s⁻¹) and quantitatively analyzed using thermal analysis (temperature range 0~900 °C, heating rate 20 °C·min⁻¹).

2.3.5. Microstructure Analysis

X-ray computed tomography (XCT, Phoenix V tome x M300, Baker Hughes, Houston, TX, USA) was utilized to reconstruct 2D and 3D microstructural morphologies and characterize pore defects of cement specimens at various corrosion ages, with scanning parameters detailed in

Table 4. XCT Scanning Parameters and Standard Settings.

Parameter	Standard Setting
X-ray tube voltage	120 kV
X-ray tube current	0.12 mA
Detector type	DXR.250
Rotation angle	360°
Detector unit	2014
Number of projections	1000
Number of pixels	2014 × 2014

. Scanning electron microscopy (SEM, SU5000, Hitachi, Japan)) was utilized for microstructural analysis, coupled with energy-dispersive X-ray spectroscopy (EDS) for elemental mapping. The samples need to be dried at 60 °C for 12 h.

3. Results

3.1. Macroscopic Morphology

Figure 2 illustrates the macroscopic morphological evolution of four cement formulations under simulated harsh formation conditions. PT specimens demonstrated the most severe degradation. Initially intact (0 d), they exhibited progressive discoloration at 7~28 d, developed macroscopic cracks at 60~90 d, and ultimately lost structural integrity, indicating the poorest corrosion resistance. The E8 group maintained uniform density throughout, showing only mild discoloration without visible damage or cracking over the 90 d period. G0.8 displayed similar morphological stability to E8. The EG formulation developed gradual discoloration during 7~60 d, followed by surface mottling and spalling at 90 d.

3.2. Compressive Strength

Under simulated harsh formation conditions, the compressive strength of the four cement groups exhibited complex nonlinear evolution over time (Figure 3a). All formulations underwent two dominant stages: short-term enhancement and long-term degradation. During early corrosion (0~7 d), all groups displayed strength increases. The EG group showed the maximum absolute strength gain, rising from 63 MPa to 73 MPa, while G0.8 exhibited the highest relative increase of 47.6% (42 MPa to 62 MPa). With prolonged exposure (7~90 d), distinct performance divergence emerged. The E8 group demonstrated superior long-term stability, stabilizing near 42 MPa after 28 d, with the strength change rate gradually decreasing from +10% (7 d) to −17% (90 d). In contrast, G0.8 showed a sharp decline from its early +45% increase to −35% at 90 d. The EG group experienced the most severe degradation, plummeting from +15% to −60% at 90 d, representing the poorest corrosion resistance among modified groups. The PT reference group failed completely by 60 d, confirming its structural instability. Notably, for all tested samples, the varying strength evolution patterns across different corrosion stages indicate no positive correlation between early-stage high strength and long-term corrosion resistance. Specifically, the EG and G0.8 groups with higher early strength growth rates show more significant strength degradation after 28 d, while the E8 group with moderate early strength growth has the best long-term strength retention.

In summary, the E8 system exhibits the highest compressive strength retention and the smallest degradation amplitude among all modified systems, which is significantly better than the PT blank group and the EG composite system. The G0.8 system is inferior to E8 but still superior to PT and EG in long-term strength stability.

3.3. Permeability

Permeability serves as a critical parameter for cement sheath integrity in CO₂ injection–production wells, being directly influenced by internal pore structure and microcrack networks. This property reflects the penetration rate of corrosive media (CO₂ and water) into the cement matrix. Typically, denser microstructures correspond to lower permeability and enhanced resistance to fluid ingress [24], thereby prolonging the service life of both cement and protected casing.

As illustrated in Figure 4a, during initial and short-term exposure (0~7 d), all groups maintained low permeability (≤0.025 mD), with E8 exhibiting optimal compactness (≤0.020 mD). After 7 d corrosion, permeability remained stable (≤0.020 mD), indicating effective pore-filling by early-stage carbonation products. Medium-term exposure (14~28 d) revealed distinct differentiation. EG and PT permeabilities increased markedly to 0.085 mD and 0.080 mD at 28 d, respectively, indicating microstructural deterioration. In contrast, E8 and G0.8 maintained low permeability (≤0.028 mD), demonstrating superior resistance. During long-term corrosion (60~90 d), inter-group differences maximized. EG permeability surged to 0.141 mD (60 d) and 0.149 mD (90 d), indicating severe structural damage. G0.8 permeability increased gradually, while E8 consistently maintained ≈0.020 mD with optimal integrity. PT specimens failed completely.

Permeability change rate analysis (Figure 4b) quantified degradation resistance. EG exhibited the most severe deterioration, rising from 372% (14 d) to 1064% (90 d). G0.8 increased steadily to 508% (90 d), outperforming PT and EG. PT showed rapid medium-term degradation (259% at 28 d), confirming inadequate CO₂ resistance. E8 maintained the lowest rate (~89% at 90 d), demonstrating superior permeability resistance and structural stability throughout the corrosion cycle.

Overall, the E8 system shows the lowest permeability and permeability growth rate, presenting a substantial improvement over the PT blank group. Among the three modified systems, E8 performs optimally, followed by G0.8, while EG displays the most severe permeability deterioration.

3.4. Porosity

Porosity variation rate serves as a reliable predictor of long-term material durability, demonstrating strong correlation with both strength retention and permeability evolution. Specimens exhibited significant differences in initial pore structure (Figure 5a). PT specimens showed the highest initial porosity (43%), followed by EG (36%), while E8 and G0.8 demonstrated superior initial compactness with 31% and 28% porosity, respectively, reflecting distinct microstructural regulation capabilities of different modifiers.

All groups displayed porosity reduction during early corrosion (0~7 d), with E8 exhibiting the most pronounced decrease (31% to 20%). However, prolonged exposure revealed divergent evolution pathways. E8 maintained optimal stability, with porosity merely recovering to 27% at 90 d. G0.8 demonstrated a clear transition from early densification to later deterioration, while EG showed the most significant rebound to 34% at 90 d. PT specimens recovered to initial porosity by 28 d.

Correlation with Figure 3 reveals an inverse relationship between initial porosity and long-term strength retention. Leveraging stable pore structure, E8 maintained 84% initial strength after 90 d corrosion, whereas EG, exhibiting the most severe pore deterioration among modified groups, retained only 39.7% strength. Notably, absolute porosity alone doesn’t determine permeability; pore connectivity and morphology are equally crucial [22]. Despite EG’s final porosity (34%) being lower than PT’s initial value (43%), its significantly higher permeability indicates formation of interconnected pore networks.

In general, the E8 system maintains the most stable pore structure with the lowest final porosity, which is remarkably better than the PT blank group. The porosity evolution further confirms that E8 has the best densification effect, followed by G0.8, whereas EG shows obvious structural degradation.

3.5. Corrosion Degree

Systematic analysis of cross-sectional corrosion morphology (Figure 6) revealed distinct corrosion gradient characteristics across all specimens. The PT group demonstrated the most rapid corrosion penetration, with substantial reduction of the central uncorroded zone by 28 d and evident corrosion-induced cracking at 60~90 d. The E8 group maintained optimal corrosion resistance throughout exposure, exhibiting slow corrosion layer development, prolonged core zone preservation, and residual spot-like uncorroded areas without cracking at 90 d. The G0.8 group displayed intermediate corrosion progression between EG and E8.

Quantitative corrosion degree data (Figure 7) further validated morphological observations. From 7 d to 28 d, PT specimens showed the most rapid corrosion advancement, reaching 73% at 28 d, significantly exceeding other groups. During this stage, E8 and G0.8 exhibited comparable corrosion degrees (approximately 55% at 28 d), indicating effective early-stage resistance. Notably, during later stages (60~90 d), the EG group experienced accelerated corrosion progression, reaching 76% at 90 d, the highest among modified groups, revealing inherent limitations in long-term durability.

A complex nonlinear relationship was observed between corrosion degree and strength retention. Despite attaining 69% corrosion degree at 90 d, the E8 group maintained 84% strength retention, attributable to its uniform corrosion mechanism that minimizes structural damage. Conversely, the EG group’s localized corrosion characteristics caused strength retention to plummet to 40% despite moderate early-stage corrosion. Similarly, permeability evolution correlated strongly with corrosion mechanism. The EG group’s maximum permeability (0.149 mD) and permeability change rate (1064%) at 90 d directly resulted from interconnected pore networks formed by structural corrosion. The E8 group’s uniform corrosion morphology prevented percolation pathway formation, maintaining minimal permeability despite a high corrosion degree. Thus, long-term performance is governed primarily by the corrosion mechanism rather than the absolute corrosion degree.

To sum up, the E8 system achieves the slowest corrosion propagation and the most uniform corrosion mode, demonstrating a significant advantage over the PT blank group. Among all modified systems, E8 exhibits the best corrosion resistance, while EG shows the worst long-term anti-carbonation performance.

3.6. XRD Analysis

To elucidate the phase evolution law of different modified cement systems during carbonation, X-ray diffraction (XRD) technology was used to systematically characterize and analyze their phase compositions (Figure 8). All diffraction patterns were normalized with the strongest diffraction peak as the reference to ensure the comparability of phase characteristics among different samples. It is worth noting that characteristic peaks of quartz (SiO₂, PDF#01-079-1910) were detected in the XRD patterns at all corrosion ages. This phase is silica fume, an additional component of the cement system, which is mainly used to improve the high-temperature resistance of cement paste, reduce the alkali content of the system, and further enhance its corrosion resistance.

In the uncorroded state (0 d), the main crystalline phase in all cement specimens was portlandite (Ca(OH)₂, PDF#00-044-1481). After curing at 130 °C, accompanied by the pozzolanic reaction of silica fume, only weak Ca(OH)₂ diffraction peaks were detected at 2θ = 47.1° and 50.8° in all modified systems, indicating that the pozzolanic reaction of silica fume effectively consumed Ca(OH)₂ in the system. In addition, the diffraction peaks of aragonite (CaCO₃, PDF#00-041-1475) appearing in the XRD patterns before corrosion were mainly attributed to the early mild carbonation reaction between trace CO₂ in the air and the cement matrix. As the main hydration product of cement, calcium silicate hydrate (C–S–H, PDF#00-034-002) gel exists in an amorphous state, so it does not show sharp diffraction peaks in the XRD patterns but presents a typical diffuse “hump peak” in the range of 2θ = 28°~30°.

After 28 d of CO₂ corrosion, distinct characteristic diffraction peaks of calcite (CaCO₃, PDF#01-072-1650) appeared at 2θ = 29.4° and 48.5° in all specimens, while the Ca(OH)₂ diffraction peak at 2θ = 50.8° disappeared completely, indicating that sufficient carbonation reaction of Ca(OH)₂ in the matrix had occurred. Among them, the intensity of the calcite diffraction peak in the EG group was significantly higher than that in the other groups, suggesting a higher degree of Ca(OH)₂ carbonation in this group. At the same time, multiple characteristic diffraction peaks of aragonite (CaCO₃, PDF#00-041-1475) were observed in the range of 2θ = 26°~53° in all systems, and characteristic diffraction peaks of vaterite (CaCO₃, PDF#01-073-1449) were detected at 2θ = 32.7° and 50.0°; the intensities of aragonite and vaterite diffraction peaks in the PT group were higher than those in the other modified groups, indicating that the carbonation decomposition of amorphous C–S–H gel in this group was more intense.

Previous studies have shown [25,26] that calcite is the direct product of the carbonation reaction of Ca(OH)₂, while aragonite and vaterite are mainly derived from the carbonation decomposition of amorphous C–S–H gel. The above phase evolution characteristics clearly indicate that during CO₂ corrosion, both Ca(OH)₂ and C–S–H gel in the cement matrix underwent carbonation reactions, and the cement system in the PT group suffered the most serious carbonation damage.

3.7. Thermogravimetric Analysis

Thermogravimetric analysis of 28 d corroded specimens elucidated modifier-specific regulation of carbonation processes. The EG group exhibited the most intense mass loss peak in the CaCO₃ decomposition range per DTG analysis (Figure 9a), indicating maximum decomposition rate, with quantitative verification showing the highest CaCO₃ content (25.21%). This correlates with the system’s high pore connectivity facilitating rapid corrosive penetration, yielding severe carbonation and the minimum strength retention (39.7%). The E8 group demonstrated moderately lower CaCO₃ content and decomposition rate, reflecting waterborne epoxy’s unique barrier mechanism. Continuous polymer films physically inhibit corrosive transport [21], confining carbonation to localized zones. The resulting CaCO₃ products preferentially fill pores rather than form connected networks, preserving structural integrity and maintaining high strength retention after 90-day exposure. G0.8 displayed intermediate carbonation behavior with moderate CaCO₃ content, where graphene’s two-dimensional nanostructure provides limited barrier effects by prolonging diffusion paths [27], yet yields suboptimal long-term performance. The PT control group showed minimal CaCO₃ content (19.92%), attributable to its rapid structural deterioration: initial high porosity enables unrestricted corrosive access, causing accelerated matrix damage and subsequent leaching of carbonation products.

4. Discussion

4.1. XCT Analysis

As established in previous studies [4,12,13,28], long-term CO₂ reaction with cement produces four characteristic corrosion zones. XCT 2D images (Figure 10) clearly reveal this stratified architecture, comprising from exterior to interior: (1) a CaCO₃ leaching zone exhibiting high grayscale values, corresponding to high-porosity regions with dissolved carbonates; (2) a dense CaCO₃ filling zone showing reduced grayscale values and substantially decreased porosity; (3) a Ca(OH)₂ depletion zone formed through decalcification of hydration products; and (4) an unreacted core zone preserving the initial matrix structure. PT and G0.8 specimens displayed distinct zonal separation, particularly pronounced in PT samples. All groups demonstrated porosity reduction in carbonated dense layers; EG’s uncorroded area developed interconnected pore networks, consistent with porosity measurements.

During early corrosion (0~7 d), CO₂ reacts with cement hydration products to form CaCO₃ and amorphous silica gel, which fill internal pores and densify the microstructure, leading to increased compressive strength in all groups. With prolonged corrosion (7~90 d), excessive carbonation decomposes C–S–H gel, destroys the cementitious skeleton, and induces microcrack propagation and pore structure deterioration, resulting in continuous compressive strength reduction.

XCT defect quantification (

Table 5. Detailed defect information of cement stone in different systems at different corrosion ages.

Sample	0 d			28 d
	Material Volume (mm3)	Defect Volume (mm3)	Defect Volume Ratio (%)	Material Volume (mm3)	Defect Volume (mm3)	Defect Volume Ratio (%)
PT	7926.38	40.13	0.50	11,023.47	605.95	5.21
E8	7641.37	18.01	0.24	10,496.90	459.90	4.20
G0.8	7346.91	24.30	0.33	10,821.84	526.66	4.64
EG	10,448.28	38.51	0.37	11,567.46	409.15	3.42

) revealed modifier-dependent initial matrix density variations. PT specimens exhibited 0.50% defect volume ratio, indicating inherent microstructural porosity. E8 achieved minimal defects (0.24%) through epoxy-mediated formation of dense organic–inorganic composites. G0.8 showed 0.33% defect volume via graphene’s barrier mechanism, while EG’s 0.37% indicated compromised microstructural optimization. Post 28 d corrosion, all groups demonstrated substantially elevated defect ratios with distinct progression patterns. PT’s defect volume surged to 5.21%, confirming rapid structural collapse. E8 increased moderately to 4.20%, demonstrating controlled corrosion inhibition. G0.8 reached 4.64%, indicating barrier failure. Crucially, EG’s lower defect volume (3.42%), coupled with its maximum permeability, suggests defect morphology dominated by interconnected pores/microcracks enabling rapid fluid transmission.

4.2. SEM Analysis

SEM analysis revealed distinct microstructural evolution across cement formulations. PT specimens transitioned from laminated to granular C–S–H morphology, maintaining porous architecture (Figure 11). BSE imaging (Figure 12) demonstrated clear corrosion stratification, with Ca EDS mapping confirming severe leaching and decalcification. These observations validate rapid structural deterioration in unmodified cement, where hydration product decalcification induces porosity escalation, fully consistent with the XCT-documented defect volume increase from 0.50% to 5.21%, reflecting the “structural degradation–macrostructural collapse” corrosion pathway.

The E8 specimens exhibited a dense and homogeneous microstructure characterized by an organic film coating, with no visible pores or cracks observed. Backscattered electron (BSE) images revealed blurred interlayer boundaries, indicating a gradual progression of corrosion. The Ca elemental mapping demonstrated optimal density and uniformity. These results confirm that the continuous organic barrier effectively retards the penetration of corrosive agents, moderates the kinetics of carbonation, and preserves the integrity of the matrix. This microstructural evidence directly corresponds to its superior macroscopic performance.

G0.8 matrices contained numerous micropores, with BSE revealing incipient layer separation and localized Ca depletion. This evidences the eventual failure of graphene’s physical barrier mechanism, where nanosheet–matrix interfaces become preferential corrosion pathways. EG specimens displayed extensive interconnected cracking networks. Though Ca mapping suggested apparent density, these permeable crack networks provide rapid transport channels, explaining the paradoxical combination of moderate defect volume with maximum permeability and minimum strength retention—revealing the microstructural essence of composite system synergistic failure.

4.3. Carbonation Resistance Mechanism of Different Modified Materials

• The early-effectiveness and late-failure behavior of the graphene-modified system originates from its physically dominated reinforcement mechanism and interfacial instability. Initially, graphene sheets reduce total porosity by filling nano-pores within C–S–H gels [29]. Well-dispersed lamellae create physical barriers, extending penetration paths through tortuous flow, thereby providing effective early-stage impermeability. The eventual failure stems from weak interfacial transition zones between graphene and hydration products [30]. Under prolonged CO₂-induced acidic and stress conditions, corrosive media preferentially attack these weak interfaces. Once debonding occurs, the exfoliated graphene sheets transform into interconnected nano-channels, causing rapid permeability increase and complete barrier failure.
• The inferior performance of the composite system stems not from individual component failure, but from a negative synergistic effect between them. It is hypothesized that the inferior performance of the composite system may be related to the thermodynamic incompatibility between waterborne epoxy resin and graphene. The hydrophobic graphene surface resists stable, uniform integration with the aqueous epoxy system. Within the cement’s porous medium, these incompatible materials undergo competitive distribution and localized phase separation. Graphene lamellae can disrupt continuous epoxy film formation, while epoxy agglomeration promotes graphene accumulation. This behavior induces micro-stresses at phase interfaces and generates inherent interconnected pores and microcracks [31]. These pre-existing connected defects create preferential pathways for corrosive media to penetrate deeply. Consequently, corrosion progresses not gradually from the surface, but directly into the interior through these channels. Rapid fluid penetration along these pathways disintegrates the cementitious skeleton, causing accelerated strength loss. Thus, although the apparent defect volume ratio may be low, defect connectivity proves substantially more detrimental than uniformly distributed closed pores.
• The superior long-term performance of waterborne epoxy resin stems from its formation of a continuous, stable, and ductile organic phase within the cement matrix, enabling regulated corrosion progression. During cement hydration, the resin forms a film and three-dimensional network structure, creating a continuous phase [23] rather than dispersed particles (Figure 13). This continuous phase establishes an effective physico-chemical diffusion barrier, significantly retarding the transport kinetics of CO₂ and H₂O. Unlike the rapid penetration in PT specimens or localized corrosion in EG, the epoxy barrier enables uniform, gradual advancement of the corrosion front. In this mode, even when surface carbonation occurs, the internal unreacted core maintains structural integrity and load-bearing capacity. The polymer confines corrosion products in situ, facilitating pore-filling rather than dissolution loss. This mechanism preserves high residual strength despite mass loss (high corrosion degree). Additionally, the polymeric film enhances cement toughness, suppressing microcrack initiation and propagation, thereby maintaining structural integrity in corrosive environments.

5. Conclusions

This study systematically evaluates the long-term carbonation resistance of four cement formulations under simulated CCUS-EOR conditions.

• Conventional Class G oil well cement (PT) fails under prolonged CO₂ exposure, exhibiting 0.080 mD permeability and 73% corrosion degree at 28 d, with complete structural failure by 60 d.
• Waterborne-epoxy-modified cement (E8) demonstrates optimal durability, retaining 84% compressive strength after 90 d. The continuous polymeric film forms a dense physico-chemical barrier, retards CO₂ penetration, maintains a uniform corrosion mode, stabilizes pore structure, and thus achieves excellent long-term performance.
• Graphene-modified cement (G0.8) shows transitional behavior. Early-stage strength increases by 47.6% via nanoscale barrier effects, but prolonged exposure induces interfacial debonding between sp²-carbon layers and the cement matrix, forming permeable nano-channels and resulting in a 508% permeability increase at 90 d.
• Waterborne epoxy–graphene-composite modified system (EG) exhibits negative synergy. It is speculated that thermodynamic incompatibility leads to phase separation and interconnected microcracks, providing fast CO₂ penetration channels, leading to serious structural damage despite moderate defect volume. It yields 0.149 mD permeability and merely 39.7% strength retention at 90 d despite moderate defect volume.
• The long-term performance of modified cement is dominated by microstructural integrity and corrosion mode rather than early strength or single corrosion index. Waterborne epoxy modification effectively balances pore structure, mechanical stability, and corrosion resistance, which is the optimal scheme for high-temperature CCUS-EOR wellbore cementing.
• In engineering practice, waterborne epoxy-modified cement is recommended for CCUS-EOR wellbore cementing in high-temperature CO₂ environments. The composite modification of graphene and waterborne epoxy is not recommended due to negative synergy, and the dosage optimization of single modifiers should be focused on in subsequent research.